Segmented copolymers with monodisperse crystallizable hard segments: Novel semi-crystalline materials

Progress in Polymer Science 36 (2011) 713–748

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Segmented copolymers with monodisperse crystallizable hard segments: Novel semi-crystalline materials Reinoud J. Gaymans ∗ Department of Science and Engineering, Twente University, PO Box 217, 7500 AE Enschede, The Netherlands

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Article history: Received 16 October 2009 Received in revised form 14 March 2010 Accepted 21 July 2010 Available online 12 August 2010 Keywords: Segmented block copolymer Monodisperse Polydisperse Nano composites

a b s t r a c t Segmented block copolymers with short monodisperse crystallizable hard segments have interesting structures and properties. In the melt, such short monodisperse segments are miscible with the matrix segments. Moreover, upon cooling, they crystallize fast, demonstrating a very high crystallinity, and only a small crystallization window is needed. The melting temperature of the short segments is high, provided that they can H-bond and/or contain aromatic groups. The melting temperature was found to decrease with increasing matrix segment concentration, due to the solvent effect of the matrix segments. At concentrations of crystallizable segment of 4–35 wt%, good dimensional and solvent stabilities were obtained. The monodisperse segments crystallized into nano-ribbons with uniform thickness and high aspect ratio, and these dispersed nano-ribbon crystallites constituted physical crosslinks, while acting also as reinforcing fillers. At concentrations of the monodisperse segments below 20 wt% no spherulitic ordering took place, and the semi-crystalline polymers were transparent. The monodisperse crystallizable segments can be used in combination with matrix segments of either low or high glass transition temperature, and may even contain (bio)functional units. © 2010 Elsevier Ltd. All rights reserved.

Contents 1.

2.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1. Miscibility . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2. Thermal mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3. Segmented block copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Monodisperse segments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Urethane . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.1. Synthesis of polyurethanes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.2. Urethanes with MDI and BDO . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.3. Other H-bonded urethanes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

714 716 716 716 717 717 717 718 718

Abbreviations: SBCP, segmented block copolymer; HS, hard segment; SS, soft segment; MD, monodisperse; PDI, polydisperse index, equal to Mw /Mn ; DII, diisocyanate; Mw , Mn , weight and number average molecular weights, respectively; Tg , crystallization temperature; Tflex , onset of rubber plateau; Tflow , flow temperature (melting temperature); Tg , glass transition temperature; Tm , melting temperature; xc , degree of crystallinity; G , storage modulus (G = 3/E); εb , elongation at break; CS, compression set; TS, tensile set. ∗ Corresponding author. Tel.: +31 534282322; fax: +31 534893823. E-mail address: [email protected] 0079-6700/$ – see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.progpolymsci.2010.07.012

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2.2.

3.

4.

Urea . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.1. Length of the urea segment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.2. PDMS–ureas . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. Diurethane–diurea . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.1. MDI . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.2. TDI . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.3. Aliphatic diisocyanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4. Urethanes with other extenders . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4.1. Diamine–diamide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4.2. LCP extenders . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5. Amide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.1. Tetra-amide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.2. Diester-amide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.3. Diester–diamide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.4. Diester–tetra-amide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.5. Non-crystallizing amide segments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.6. Type of the matrix segment A . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.6.1. Amorphous segments with a low Tg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.6.2. Semi-crystalline segments of low Tg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.6.3. Extended soft segments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.6.4. Amorphous matrix segments with a high Tg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.6.5. Semi-crystalline matrix segments with a high Tg and high Tm . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.7. Other copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Polydispersity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. Melting temperature of HS (DSC and DMA) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4. Crystallinity (DSC and FTIR) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5. Rate of crystallization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.6. Thermal stability . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7. Morphology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7.1. TEM . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7.2. AFM . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7.3. WAXS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7.4. SAXS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8. Dynamic mechanical properties (DMA) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.1. Tg soft segment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.2. Soft Segment crystalline phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.3. Modulus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.4. Temperature dependence of the rubber plateau modulus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.5. Hard segment Tg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.6. Flow temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8.7. G loss modulus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.9. Melt rheology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.10. Tensile behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.10.1. Young’s modulus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.10.2. Yield behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.10.3. Past the yield point . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.10.4. Structural changes upon drawing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.10.5. Fracture behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.11. Elastic properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.11.1. Compression set . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.11.2. Tensile set . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.11.3. Stress relaxation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.12. Fiber properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.13. Solvent resistance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.14. Hydrophilic copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction Synthetic polymers are still far from the elegant schemes of biological system in self-organization phenom-

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ena and formation of highly complex tertiary structures with specific functions [1]. Over the past 40 years, segmented block copolymers (multi-block copolymers) with short monodisperse crystallizable segments have been

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Fig. 1. Segmented block copolymers with either mono- or polydisperse crystallizable hard segments and their crystalline packing.

studied, starting with the work of Harrell [2] and Ng et al. [3]. Segmented block copolymers (SBCPs) are an intriguing type of copolymers. They consist of linear copolymers with alternating A- and B-segments of telechelic units, where the B-segments often crystallize and demonstrate high melting temperatures (Fig. 1). These segmented block copolymers are semi-crystalline materials in which the short B-segments may crystallize in their extended form. If the segments are monodisperse (MD) in length, they stack neatly and form crystals of uniform thickness, without a partial crystalline interphase layer that is present in copolymers with polydisperse B-segments (Fig. 1). Copolymers with these monodisperse segments have well-defined crystallite structures in which the crystallites are dispersed in a matrix phase of the A-segments. In such well-defined copolymer systems, structure–property relationships can be more easily studied. Also, these semi-crystalline materials present a very interesting combination of physical and engineering properties. Depending on the segments that form the matrix phase, several types of materials are obtained. The typical thermomechanical behavior of four types of materials is given in Fig. 2.

Fig. 2. The DMA behavior of segmented block copolymers with monodisperse, high melting crystalline segments (∼10 wt%) and a matrix phase consisting of segments that are: a, amorphous and of low Tg ; b, amorphous and of high Tg ; c, semi-crystalline and of low Tg as well as low Tm ; d, semi-crystalline and of higher Tg and high Tm .

In their simplest form, these copolymers present a twophase morphology in the solid state: an amorphous matrix phase consisting mainly of A-segments and a dispersed phase of B-crystallites (Fig. 2a and b). For temperatures between their Tg and the Tm , these copolymers display a low modulus and a high elasticity. The dispersed high melting crystallites represent the physical crosslinks of the matrix segments and provide the copolymer with dimensional stability and solvent resistance [4–9]. The high melting B-segments are often called the hard segments (HS), and when they are monodisperse, they are denoted as MD-HS. If the A-segments also crystallize to a certain extent, a three-phase morphology is obtained: an amorphous Aphase, a crystalline A-phase and a crystalline B-phase of elevated melting temperature (Fig. 2c). Due to the presence of the crystalline B-phase in the copolymers, the semicrystalline A-phase can be well oriented at a temperature very near its melting temperature. During the orientation of the crystallizable A-phase, strong strain-induced crystallization of that phase can take place. With such a strain-induced crystallization the fracture stresses are strongly increased. The A-phase can also have a high Tg and a high Tm , like a polyester (Fig. 2d). These block copolymers can be highly oriented at a temperature very near the melting temperature of the A-segments, thereby increasing the orientation and the fracture stress. The key to these copolymers is the short monodisperse crystallizable B-segments that, in the melt, are miscible with the A-segments so that high-molecular-weight copolymers can be readily synthesized. Upon cooling, these B-segments crystallize very fast, and to a large extent, into crystallites with a uniform thickness. Important issues for the B-segments include:

-

the miscibility with the A-segments in the melt; their rate of crystallization upon cooling; their degree of crystallinity; their melting temperature; their stability against trans- and degradation-reactions; the morphology of the B-segment crystallites; how well the B-segment crystallites resist deformation; and how well the B-segment crystallites resist dissolution.

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The present review discusses the synthesis, structure and properties of these materials.

1.1. Miscibility The miscibility of segmented multi-block copolymers has little been modeled [10,11]. The miscibility of the Aand B-segments is often limited and in cases where they are immiscible, liquid–liquid demixing takes place. The miscibility of the A- and B-segments is governed by the compatibility of the structures, e.g., as reflected in their Flory interaction parameters, . The miscibility is also affected by the A- and B-block lengths and A- and B-block polydispersity, but only slightly by the total molecular weight of the copolymer [10–13]. With increasing polydispersity the miscibility decreases due to the increase in the Flory interaction parameter  [10]. If either the A- or B-segments are longer than a certain length, micro-phase separation occurs [11,13,14]. The miscibility in polyurethanes increases with temperature. In some situations on cooling from a homogenous melt, liquid–liquid demixing takes place. However, often in polyurethanes the demixing process is slow and a sample cooled fast from a homogenous melt will remain homogenous at low temperatures, with phase separation taking place on heating, followed by rehomogenization at high temperatures [15]. However, with longer A- and/or Bsegments the melt is often already (partially) liquid–liquid demixed. Particularly highly polar B-segments (such as urethane, urea, amide) are poorly miscible with apolar Asegments, and the miscibility is very sensitive to the length of the B-segments and the number of H-bonding units. If the HS structure is regular, then on cooling crystallization can take place. If both liquid–liquid demixing and crystallization of the B-segments occur, the morphology becomes complex and the properties are not optimal [14]. If liquid–liquid demixing is unwanted, there is a dilemma: the B-segments must be long enough to crystallize and have a high melting temperature, but not so long that they exhibit liquid–liquid demixing. For crystallization, the minimal average length of the polydisperse HS is often three or four repeat units [14]. Liquid–liquid demixing with polar B-segments can take place already with a length of two to three repeat units [14]. The margins for crystallization without liquid–liquid demixing are thus often small. In polydisperse systems, there is a segment length distribution, and both short and long B-segments are present. The very short B-segments have difficulties crystallizing and the long B-segments demix. One can imagine that if the B-segments are monodisperse in length, their length can be tuned in order to be sufficiently short that liquid–liquid demixing is prevented but long enough to crystallize upon cooling giving rise to a high melting temperature. Copolymers with short monodisperse B-segments have the advantage of not undergoing liquid–liquid demixing, still being able to fully crystallize and have crystallites with a uniform thickness. For this reason, segmented block copolymers with short monodisperse crystallizable segments have been extensively studied.

Fig. 3. The thermal mechanical behavior of –(PMTO2016 –MDI–(BDO–MDI)3 )n – of monodisperse (solid line) and polydisperse (dashed line) urethane segments. Reproduced with permission from ref. [8], copyright (1986) Springer–Plenum.

1.2. Thermal mechanical properties The SBCPs with MD-HS that have been studied have mostly been based on polyether matrixes and urethane, urea or amide crystalline segments. The MD-HS components were short, on the order of 1–5 repeat units, and their concentration in the copolymers was low (5–35 wt%). The difference between SBCPs with monodisperse and random disperse B-segments is clearly illustrated in the thermal mechanical behavior of certain polyurethanes (Fig. 3). The SBCPs with monodisperse B-segments demonstrate a lower and sharper Tg , a rubber plateau where the modulus is only slightly dependent on the temperature, and a sharper melting transition. The Tg of the A-segment was low, since little HS was dissolved. The Tg transition was sharp due to the absence of liquid–liquid demixed amorphous B-segments. Moreover, the modulus in the region between Tg and Tm was little dependent on temperature as the short MD-HS crystals melted at a single temperature. The modulus in the region between Tg and Tm was relatively high as the crystallinity of the MD B-segments was high. 1.3. Segmented block copolymers Segmented block copolymers are generally linear multiblock materials of alternating A- and B-segments, often consisting of some 20–50 segments. The high melting Bsegments constitute the physical crosslinks and reinforcing fillers for the A-segment phase and give the material its dimensional stability and solvent resistance. These materials can be easily melt processed as well as reprocessed. The segmented block copolymers are synthesized starting from telechelic A-polymers that have functional groups such as hydroxyl, amine or acid moieties. The B-units can be prepared with functional end-groups (isocyanate, acids, esters) prior to the polymerization or in situ during the synthesis. Examples of monodisperse crystallizable Bsegments that have been used include urethanes, ureas, amides, esters and combinations of these. The urethane, urea and amide groups can form H-bonds with their neigh-

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Fig. 6. Synthesis of PTMO, MDI and BDO polyurethanes by: a, a one-pot method; b, a two-step method; c, an alternative two-step method.

Fig. 4. Possible H-bonding with urethane, urea and amide groups.

bors (Fig. 4) and the H-bonding strength increased in the order of urethane < amide < urea. With increasing Hbonding strength the melting temperatures increase and the crystallites are more resistant to deformation and solvents. Moreover as non-crystallizing matrix segments with a low Tg are often used; poly(propylene oxide) (PPO) and poly(dimethyl siloxane) (PDMS). These SBCPs are thermoplastic elastomers with excellent low-temperature properties (Fig. 2a). Amorphous engineering polymers, that have a high Tg , e.g., poly(phenylene oxide) (PPE), polysulphone (PSU) and polycarbonate (PC), have also been copolymerized with monodisperse crystallizable Bsegments (Fig. 2b). The solvent resistance of these high Tg polymers was strongly improved with these B-segments, even at low concentrations. The A-segments, the matrix material, can also be semicrystalline, i.e., with a Tm in addition to a Tg (Fig. 2c). Typical examples of semi-crystalline A-segments include poly(tetramethylene oxide) (PTMO), poly(ethylene oxide) (PEO), poly(hexamethylene adipate) (PHMA) and polycaprolactone (PCl). With these semi-crystalline matrix segments, strain-induced crystallization can take place and high-strength thermoplastic elastomers are obtained. Copolymers with matrix segments that have both high Tg and Tm (Fig. 2d), such as poly(methylene terephthalate)s, have as of yet only been studied to a very limited extent. Also, attention has been paid to di- and tri-block copolymers, and particularly their melt rheological behavior was found to be interesting. 2. Monodisperse segments In segmented block copolymers, both the matrix segments A and the phase-separated segments B can be monodisperse (MD) in length. If the A-segments are MD

in length, this had only a small effect on their glass transition temperature, on the phase separation of the B-segments and on the properties of these copolymers in general [3,8,9,16–20]. However, monodisperse crystallizing B-segments have a significant influence on the morphology and properties of the copolymers. 2.1. Urethane The first studies on monodisperse (MD) segments involved MDI (4,4 -methylenebis(phenyl isocyanate) and the non-H-bonded piperazine (Pip)-based polyurethanes (Fig. 5) [2,3,8,9,12,21]. Unfortunately, the MDI urethanes are thermally unstable above 180 ◦ C, which limits the melt synthesis and the melt processing of these copolymers [8]. However, the nonH-bonded piperazine urethanes present a higher thermal stability as the piperazine segments are less polar. Piperazine PUs have mostly been reported on together with PTMO as the matrix segment [2,3,9,21]. 2.1.1. Synthesis of polyurethanes PUs can be prepared in several ways (Fig. 6). In the one-pot method, the polyether prepolymer, the diisocyanate (DII) and the butane diol (BDO) extender are mixed, (Fig. 6a) and the formed HS demonstrates a polydisperse length distribution. One should bear in mind that the reactivity of the two isocyanate units of a DII often differs [12,22,23]. The first isocyanate group in the MDI is somewhat more reactive than the second and this makes the HS distribution more narrow than truly random [12,22,23]. For a polyurethane –(PTMO1000 –MDI–(BDO–MDI)x )n – with x = 1, obtained by the one-shot method, a polydispersity index (PDI) of 1.67 has been measured [12]. The –(PTMO1000 –MDI–(BDO–MDI)1 )n – copolymer can also be synthesized in the two-step, one-pot method described in Fig. 6b. In this procedure, the A-segments are first endcapped with MDI groups, before the chains are extended with BDO units. If the MDI endcapped prepolymers are not purified from the unreacted MDI units, the HS distribution in the copolymer becomes narrow giving rise to a PDI for the urethane unit of 1.37 [12]. With the

Fig. 5. Chemical structures of –MDI–(BDO–MDI)x – H-bonded urethane and –Pip–(BDO–Pip)x – non-H-bonded urethane.

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narrow-disperse HS, as compared with the polydisperse HS, the SS glass transition is sharper, the plateau modulus less temperature-dependent and the flow temperature higher. Also the elastic properties are improved, with lower tensile set values and hysteresis energy. The two-step method using an excess of the DII in combination with washing out of the excess has also been used. The intermediate DII–SS–DII is more pure and can be extended with a diol (or diamine). In this way copolymers with –(PTMO–MDI–BDO–MDI)n – copolymer with monodisperse tetra-urethane segments –MDI–BDO–MDI– have been prepared [12]. If a longer MD PU segment is wanted, the diol extender unit can be a diol–diurethane, a diol–tetra-urethane, etc., prepared separately. Copolymers with monodisperse PU segments can also be prepared by first synthesizing a monodisperse diisocyanate endcapped unit and reacting it with the dihydroxy SS according to Fig. 6c. Yet another method is to start with a SS-2OH and stepwise increase the urethane group length with purification after each step and finishing off with a SS-2OH [8]. This way, MD segment lengths as long as 6 repeat units have been prepared. 2.1.2. Urethanes with MDI and BDO Eisenbach et al. [8] and Christenson et al. [20] have studied the influence of the length of monodisperse PU oligomeric units on the lamellar thickness and the melting temperature. The investigation involved MDI–(BDO–MDI)x model compounds, (MDI–BDO)x MDI–PTMO2016 –MDI–(BDO–MDI)x tri-block copolymers and –PTMO2016 –MDI–(BDO–MDI)x – multiblock copolymers. The repeat length x of the monodisperse units was varied from 0 to 5, and the MDI-based MD-PUs were synthesized at low temperatures in solution to prevent trans- and degradation-reactions from taking place [8,24]. For similar reasons, the mechanical properties were studied on solution cast films. The crystallite thickness of the urethanes, as determined by WAXS, increased linearly with the length of the PU unit and, at the same time, the melting temperature was raised. The latter is known to increase linearly with the (reciprocal) crystallite thickness, and a general accepted model for lamellar crystals is that of Hofmann and Weeks [25]. Blending of monodisperse urethanes of varying lengths did not lead to the formation of mixed crystals, but to eutectic-type phase diagrams which typified the polymorphic character of these mixed systems [8]. Upon heating the monodisperse PU units to above their melting temperature, a polymorphic system was obtained and this was due to the fast trans-urethane reactions in the melt [8,26]. The IR spectra of these PUs presented crystalline carbonyl bands at 1690–1710 cm−1 and an amorphous carbonyl band at 1735 cm−1 . The crystalline carbonyl band position and size were dependent on the thermal history whereby the solution cast and the annealed urethanes gave the lowest wavenumbers and the highest intensities. As compared to PUs with a random length distribution, the monodisperse polyurethanes had strikingly different thermal and thermo-mechanical behaviors, particularly with regard to the plateau modulus and the flow temperature

Fig. 7. The thermo-mechanical behavior of monodisperse –(PTMO2016 –MDI–(BDO–MDI)x )n – polyurethanes when varying the urethane segment length x. Reproduced with permission from ref. [8], copyright (1986) Springer–Plenum.

(Fig. 3). Increasing the urethane segment length x from 1 to 5 led to a significant increase in the modulus of the plateau at room temperature as well as of the flow temperature (Tm ) (Fig. 7). Also the plateau modulus was little dependent on temperature. However, the onset of the rubber plateau shifted to higher temperatures, albeit not as much as with the polydisperse systems [8]. 2.1.3. Other H-bonded urethanes MDI-based polyurethanes are the most often studied. Polyurethanes with monodisperse HS can also be obtained from other DII, such as toluene diisocyanate (TDI) [27–29], p-phenylene diisocyanate (p-PDI) [28–30], m-phenylene diisocyanate (m-PDI) [28–30], hexane diisocyanate (HDI) [28–31], butane diisocyanate (BDI) [18,26,31–34], hydrogenated MDI (HMDI) [28–30], and cyclohexane diisocyanate (CHDI) [28–30] (Fig. 8). If the chemical structure of the DII is regular (symmetrical), as for p-PDI, BDI, HDI, the ease of crystallization is

Fig. 8. Structures of various diisocyanates.

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Fig. 9. Synthesis of copolymers with diurethane segments.

improved, and the crystallinity and melting temperature are higher. 2.1.3.1. Diurethane segments. Materials that are very simple to prepare and with a well-defined structure are copolymers with diurethane segments [28–30]. Yilgor and co-workers [28–30,35,36] studied copolymers based on PTMO975 and diurethane HS by reacting the PTMO with the DII at a ratio of 1:1 (Fig. 9). The segments are very short and automatically monodisperse in length. The concentration of these very short diurethane segments with PTMO1000 SS is low, 14–20 wt% (Table 1). If the urethane structure is irregular, as for 2,4-TDI and m-PDI, two transitions for the copolymer become apparent: a Tg of the PTMO phase (at −60 ◦ C) and a Tflow of the diurethane segments, just above room temperature. An SBCP with a PTMO1840 –TDI based diurethane was found to be a viscous liquid at room temperature and no micro phase separation of the TDI groups was noted [28]. If, on the other hand, the diurethane had a more regular (symmetrical) structure, the Tflow occurred at a higher temperature. This led to the MDI-based diurethane presenting two transitions; a Tg at −60 ◦ C and a Tflow at 120 ◦ C. Nevertheless, the copolymers remained very soft materials [10,30]. The morphology of the MDI PU, as studied by AFM, did not show crystalline ribbons and no peaks were seen with SAXS. The copolymer did, however, have a very high fracture strain (1600%). If the regular p-PDI or 1,4-cyclohexane diisocyanate were used, the copolymers had more than two transitions: a Tg at −70 ◦ C, a transition at about 60 ◦ C and a transition above 100 ◦ C. The latter two transitions must have come from the diurethane units. Moreover, at 50 ◦ C, the p-PDI crystalline phase had almost disappeared [30]. With DSC, a melting temperature at 50 ◦ C could also be observed [30,35,36]. For CHDI copolymers crystalline ribbons were observed with AFM and crystalline peaks in the SAXS spectrum [37]. The p-PDI and CHDI diurethane copolymers displayed E-moduli at approximately 30–50 MPa, elevated fracture strains and not too high tensile set values. The higher mod-

ulus of the CHDI copolymers might have been due to a higher aspect ratio of the crystalline ribbons. For the p-PDI copolymer, the crystallization in time at room temperature was monitored with FTIR [28,29,35,36]. With time, the 1695 cm−1 peak was found to increase to the detriment of the 1733 cm−1 peak height while the 1708 cm−1 peak remained fairly constant [35]. At the same time, the Estorage modulus increased during 24 h from 30 to 90 MPa. Also the aliphatic HDI units in diurethane were found to crystallize well, and with a Tflow of 45 ◦ C. This copolymer also had an E-modulus of 23 MPa, a very high elongation at break (1380%) [36]. The decrease in modulus of these diurethane copolymers over the temperature range −30 to 50 ◦ C was less than one decade [30]. 2.1.3.2. Tetra-urethane segments. Polyurethanes with tetra-urethane segments are generally prepared with butane diol as the PUs with this diol present a high crystallinity and high melting temperature. Compared to the diurethane segments, the tetraurethane units have longer segment lengths, more H-bonding units and a higher concentration in the copolymer. Consequently, the copolymers display a higher modulus and a higher melting temperature. With the irregular TDI the copolymer was at room temperature still a tacky solid. The melting temperature with the –BDI–BDO–BDI– segments changed with the type of SS and the HS concentration (Table 2). These copolymers also demonstrate high fracture strains (∼800%), albeit lower than those reported for the diurethane copolymers (950–1600%) (Table 1). 2.1.3.3. Other monodisperse polyurethane systems studied. 2.1.3.3.1. TDI–(BDO–TDI)m –HS. TDI is an extensively used diisocyanate due to its low price. It is often sold as a 80/20 mixture of 2,4- and 2,6-TDI. With 2,4-TDI, the chain regularity is poorer than for MDI and this is even more the case with the mixture of 2,4- and 2,6-TDI. Polyurethanes with monodisperse TDI2,4 –BDO-based HS have been studied (–PTMO1000 –TDI–(BDO–TDI)x –) and the number of repeat units x was changed from 0 to 6 [27,38]. These –TDI–(BDO–TDI)x – segments were

Table 1 Some properties of –(PTMO1000 –DII)n – copolymers with diurethane segments. DII

MDI 2,4-TDI m-PDI p-PDI CHDI HDI a b

HS (wt%)

20 15 14 14 14 15

Tg (◦ C)

−60 −60 −60 −70 −70 −70

The G-modulus corresponds to E-modulus/3. Tacky solids.

Modulus (MPa) E

G a

7 2 5 29 50 25

2.3 0.7 1.7 10 17 8.3

Tflow (◦ C)

εb (%)

Ref.

120 35 35 60/125 75/>180 45

1600 –b –b 950 1150 1400

[12,30] [30,35,36] [30,35,36] [30,35,36] [30,35,36] [30,35,36]

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Table 2 Various properties of tetra-urethane –(SS–DII–DBO–DII)n – copolymers. DII

SS (g/mol)

HS (wt%)

Tg (◦ C)

Modulus (MPa) E

MDI 2,4-TDI BDI BDI a

PTMO1000 PTMO1000 PTMO2000 CL1000

39 32 15 25

−50 to −42 −35 −45 –

Tflow (◦ C)

εb (%)

Ref.

95–150 20 70 105

– 800 850

[8,12] [38] [32] [17]

a

G

30 10 Tacky solid 23 8 – –

The G-modulus corresponds to E-modulus/3.

not crystalline, despite that they were monodisperse in length. In addition to a Tg from the PTMO, a separate Tg of the urethane segments was present from x = 3 and its temperature increased with the segment length. At x = 3, it was −18 ◦ C as opposed to 70 ◦ C for x = 6. Thus at x > 2 the –TDI–(BDO–TDI)m – segments were phase-separated and were probably present as dispersed nano-particles. When increasing the MD-HS length (content), the fracture stress increased significantly; at a repeat segment length (x) of 6, it had reached a value of 8 MPa. The fracture strains were also elevated, e.g., 1100–2100%. SAXS studies on these copolymers have revealed a weak scattering, suggesting a poor domain organization [38]. IR-analysis has indicated that the N–H groups were fully H-bonded whereas the H-bonding of the carbonyl of the urethane group was only half [39]. Part of the H-bonding of the N–H groups was believed to occur with the ether groups. Furthermore, the concentration of TDI dissolved in the SS phase was high. Also TDI–BDO-based segmented block copolymers with polybutadiene (Mn 2200 g/mol) and polycaprolactone (Mn 1630 g/mol) SS have been studied [27]. When using PB2200 , the Tg of the SS was somewhat lower (∼−60 ◦ C) and the Tg of the –TDI–(BDO–TDI)m – was apparent already at x = 2. The polar urethane and the non-polar polybutadiene segments phase-separated easily. When using PCL1640 , the DSC traces were more complex as a result of the PCL being a semicrystalline segment with a Tm of ∼50 ◦ C. Monodisperse PUs with the more regular 2,6-TDI segments have also been investigated and the Tg s of the PU in these copolymers were higher than for the similar 2,4TDI–BDO [27]. Thermal mechanical analysis revealed that, in addition to the Tg of PU as observed with DSC, other transitions occurred at higher temperatures. These transitions were sensitive to the processing conditions [38]. 2.1.3.3.2. BDI–BDO. For biomedical applications, copolymers with monodisperse aliphatic polyurethanes based on BDI and BDO have been studied, due to their degradation products being thought to be unharmful [32]. The MD urethane segments consisted of –BDI–(BDO–BDI)x –, where x was 1 or 2. As SS, a polycaprolactone with a molecular weight of 2000 g/mol was selected. When going from –BDI–BDO–BDI– to –BDI–(BDO–BDI)2 – the Young modulus increased from 23 to 70 MPa and the HS melting temperature from 70 to 140 ◦ C. 2.1.3.3.3. Non-H-bonded urethane. The non-H-bonded monodisperse PU HS–(Pyp–(BDO–Pyp)x )– were the first segmented block copolymers with MD segments to be studied in detail. They consist of piperazine-based

polyurethane groups according to Fig. 5 [2,3,9,21]. An advantage of these non-H-bonded PUs is that they are more thermally stable, thus preventing the monodisperse structure from being randomized at elevated temperatures [9,21]. The higher thermal stability of the non-H-bonded polyurethanes is due to the absence of active hydrogens in the urethane groups [21]. An additional advantage of the non-H-bonded urethane is its lower polarity resulting in superior phase mixing with the PTMO segments, and thus a suppression of liquid–liquid demixing. As a result, even long non-H-bonded urethane segments remained miscible in the melt. Harrell [2] was the first to study these piperazine urethanes, and later, the same materials were also investigated by Ng et al. [3], Wilkes and Samuels [40] and Eisenbach et al. [1,9,21,41]. The synthesis was carried out in steps: First, the urethane units were synthesized and coupled to the PTMO prepolymers, according to Fig. 6c. The repeat length x of the urethane units –Pyp–(BDO–Pyp)x – was varied from 1 to 4. A certain difficulty was encountered in the purification of the higher members of the urethane units and, as a result, the distribution of the PU segment length was narrow. Also copolymers with a polydisperse distribution were prepared and the properties compared [2,3]. The copolymers with MD-HS had a nano-ribbon morphology which was very stable, as evidenced by a 30-year-old sample having maintained its morphology [42]. Eisenbach and co-workers [1,9,21,41] performed detailed studies on piperanzine PUs with monodisperse segment lengths. The PU precursors were prepared by a step-by-step method with intermediate recrystallization. The monodisperse PU units presented lengths of 1–7 repeat units (x) and their purities were above 99%. The copolymers based on these precursors and PTMO2000 were of a high molecular weight. Some model compounds of these PU units were also studied. In the crystalline state, the urethane model compounds had a fully extended conformation, as observed by WAXS, and the melting temperatures increased with the unit length. Mixtures of model compounds were found to have a cocrystallization behavior when the repeat length x was at least 4 and if the unit length difference was only one repeat unit. In other circumstances, an eutectic point was observed in the composition graph. The melting temperatures in the PU copolymers also increased with the HS segment length, but were slightly lower than those of the model compounds. The PTMO SS had a solvent effect on the melting temperature of the HS. A mixture of copolymers with varying PU segment lengths showed multiple transitions, similarly to the

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Fig. 10. The synthesis of copolymers with diurea segments: (a) with a DII; (b) with dimethyl carbonate and a diamine.

model compounds [9,41]. Copolymers with polydisperse HS also presented multiple transitions, and after annealing, the crystals became reordered by hard segment length. The heat of fusion for short HS was generally lower in the copolymer than in the model compound, which might be the result of differences in crystallization and/or due to the effect of the SS on the heat of fusion. However, with long HS, the results of the copolymer and the model compounds were comparable. The thermal mechanical behavior of these copolymers was analyzed by DMA. The trends observed on –(PTMO–(Pyp–(BDO–Pyp)x )n – were quite similar to those on –(PTMO–MDI–(BDO–MDI)x )n – (Figs. 3 and 7). It was quite striking that the temperature dependence of the modulus between the room temperature and the HS Tm was so small [9,41]. This high dimensional stability suggests that all HS crystals melted over a narrow range. A comparison of the MD-HS copolymers with their random length HS counterparts revealed that the transition temperatures now hardly changed. However, for the MD-HS, the room-temperature modulus was a factor 2.5 higher. The crystallizable monodisperse non-H-bonded urethane HS was thus present in the crystallites in their extended form [1,9,43]. For the introduction of a unit incapable of extending, e.g., an o-benzoic group, chain folding took place [43]. The folded crystals had half the thickness of their extended counterparts and as a result of this chain folding, the melting temperature was 30 ◦ C lower. The non-H-bonded polyurethanes thus had a comparable crystallization behavior to that of the H-bonded PUs. 2.2. Urea Ureas are interesting segments as they can form next to mono-dendated also very strong bidentated H-bonds (Fig. 4). The H-bonding is stronger than for polyurethanes and polyamides. When a diisocyanate molecule is placed between amine-terminated soft segments, a monodisperse diurea HS is formed (Fig. 10a). The urea groups can also be formed in two steps: first the diamine SS end-groups react with excess dimethyl carbonate to give a diurethane unit (Fig. 10b). Such an SS with a diurethane unit can react with diamineterminated soft segments, leading to copolymer with diurea blocks. Diurea copolymers with amine-terminated PTMO [26,30,31,34–36], PCl [28,44] and PDMS [45–48]

have been studied. The influence of the DII on the diurea copolymers with ∼PTMO1000 as SS was investigated with several DIIs (Table 3). The diurea units are short in length, and so had for example the p-PDI based diurea an extended length of only 1.0 nm. The melting temperature of the diurea copolymers depends on the structural regularity as well as the stiffness of the diurea group [30,31]. Particular diurea segments with a high structural regularity (symmetry) and rigidity, such as p-PDI and the CHDI, have displayed remarkably high moduli and melting temperatures (Tm > 250 ◦ C). There is also an effect of the type and length of the soft segment [30,31]. Due to their strong H-bonding, diurea copolymers have much higher melting temperatures than copolymers with diurethane groups (+100 ◦ C) (Tables 1 and 3) [36]. Moreover, the diurea copolymers present higher moduli, but lower fracture strains [30]. For all of these copolymers, the modulus in the region between room temperature and the melting temperature has been found to be very little dependent on temperature [30]. Morphological studies by AFM have presented ribbon-like (crystalline) hard domains with high aspect ratios (Fig. 11) [26,30,34–36]. SAXS analysis has been performed on copolymers comprising diurea segments with extended lengths in the order of 1 nm. The observed interdomain spacings for p-PDI, m-PDI, CHDI and HDI were respectively 5.5, 5.5, 7.9 and 7.9 nm [36]. For BDI, a long spacing of 6.0 nm as well as an interdomain spacing of 4.5 nm has been reported [26]. These are remarkable differences considering that the diurea unit length and SS length hardly changed. The monodisperse aliphatic diurea segments have been found to crystallize extremely fast [26]. Decreasing the methylene length of aliphatic diurea HS led to an increase in the melting temperature, however not for the very short ethylene group (Table 3) [31]. Furthermore, with an even number of methylene units, as compared to an uneven number, higher melting temperatures were obtained. When ethylene units were present, the melting temperature was lower (115 ◦ C) than expected. Overall, the aliphatic diurea HS were found to provide decent properties. The use of BDI is interesting for biomaterial applications as its degradation product is the tetramethylene diamine (putrescine), a naturally occurring diamine in the body. An HDI-based diurea copolymer showed no signs of degradation when tested for 30 min at 170 ◦ C [31]. Among the various diurea copolymers, aliphatic diisocyanates,

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Table 3 Various properties of –(PTMO1000 –diurea)n – copolymers. DII

MDI HMDI 2,4-TDI m-PDI p-PDI CHDI EDI BDI HDI a b

HS (wt%)

20 21 15 13 13 13 10 12 14

Tg (◦ C)

−60 −60 −60 −60 −60 – – −50 −60

Modulus (MPa)

εb (%)

Ref.

∼190 ∼130 52 130 >250 >250 115 140 150/126

700 1200 800 700 550 800 – 1060 990/800

[30] [30] [30] [30,35,36] [30,35,36] [30,36] [31] [26,31] [30,31,36]

a

E

G

46 6

15 2 –b

63 76 ∼70

21 25 ∼23 –

96 94/120

Tflow (◦ C)

32 31/40

The G-modulus corresponds to E-modulus/3. Tacky solids.

such as HDI and BDI, gave systems with melting temperatures below 200 ◦ C. This suggests that these copolymers can be melt processed without severe degradation. 2.2.1. Length of the urea segment The influence of increasing the number of urea groups in the HS segments was studied on butane diisocyanate BDI copolymers –(PTMO–Ur–((CH2 )4 –Ur)x )n – [31]. Copolymers with mono- to tetra-urea groups (x = 0–3) in the segment were investigated. An FTIR analysis on these copolymers focused on the position of carbonyl band, which is sensitive to the H-bonding packing. For the case of the mono-urea, the copolymer was in the melt phase at room temperature and a urea carbonyl peak could be observed at 1637 cm−1 . For the copolymers with di-, tri-, and tetra-urea segments, the urea carbonyl peak occurred at 1615 cm−1 and a small peak was present at 1701 cm−1 . Upon warming a sample with two urea groups (x = 1), the

1615 cm−1 peak decreased in size while those at 1635 and 1690 cm−1 increased, both for the mono-dendated Hbonding and the non-crystalline urea [26]. Upon heating to above 200 ◦ C, also the 1640 and 1690 cm−1 peaks disappeared at the cost of a broad 1730 cm−1 peak. This change was irreversible and was probably due to degradation of the urea groups at such high temperatures. A random disperse diurea sample had a carbonyl peak at slightly higher wavenumbers (1620 cm−1 ) [26]. This indicates that, on average, the H-bonding in the polydisperse HS was somewhat weaker. The flow temperatures of the copolymers increased significantly with the number of urea groups in the HS: from <20 ◦ C for the mono urea to well above 200 ◦ C for the materials with tri- and tetra-urea groups. The difference in flow temperature on a monodisperse and a polydisperse diurea copolymer was small (140 ◦ C as compared to 135 ◦ C). 2.2.2. PDMS–ureas Polyureas with poly(dimethyl siloxane) (PDMS) soft segments, for which the molecular weights range from 900 to 3660 g/mol, have been studied [45–48]. The diurea copolymers were synthesized from amine-functionalized poly(dimethyl siloxane) and a diisocyanate (Fig. 10a), and the diurea group was based on TDI, MDI and HMDI. As PDMS is highly apolar (ı ∼ 7.5 (cal/cm3 )1/2 ) and the MDI urea is very polar (ı ∼ 11.3 (cal/cm3 )1/2 ), a solvent was used to carry out the synthesis [45]. Three transitions were observed in DMA graphs of the PDMS–MDI copolymer, at −100, 65 ◦ C and a higher temperature. The first two transitions corresponded to the Tg of the PDMS and the MDI urea segments, respectively. 2.3. Diurethane–diurea

Fig. 11. An AFM phase angle mode image of PTMO2000 –BDI diurea. Reproduced with permission from ref. [26], copyright (2005) American Chemical Society.

As mentioned above, copolymers with diurethane segments display rather low melting temperatures. For tetra-urethane copolymers, the melting temperatures are somewhat higher but still outside the practical range. The diurea copolymers have sufficiently high melting temperatures combined with relatively elevated moduli. However, these diurea copolymers are prepared from amine-functionalized polyethers, which are less available than hydroxyl-functionalized prepolymers. Copolymers with diurethane–diurea segments may be prepared by

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723

Fig. 12. The synthesis of copolymers with diurethane–diurea segments: (a) a DII-endcapped prepolymer; (b) a diurethane–diurea copolymer.

first endcapping the dihydroxy polyether with diisocyanate groups and then extending these with a diamine (Fig. 12) [7,16,30,49–54]. This synthesis is simple; however, if the diisocyanateendcapped prepolymer is not purified from unreacted DII, the hard segment (HS) distribution in the copolymer will be narrow. 2.3.1. MDI Well-known examples of polyurethane–urea include the Spandex materials (e.g., Lycra) based on PTMO2000 , MDI and ethylene diamine. The effect of the diamine length has been studied [7,16,50] in the system of PTMO2000 /MDI/diamine, where it was varied from 0 (hydrazine) to 6 (hexamethylene diamine, HMDA). All the resultant copolymers presented soft segment (SS) Tg s (∼−30 ◦ C) with very little dependent on the diamine length. The melting temperatures of these narrowly dispersed HS increased slightly with a decreasing diamine length (150–180 ◦ C). Also, the copolymers with an even number of diamines had higher melting temperature. The plateau modulus increased with increasing diamine length and was nearly independent of temperature [16,50]. 2.3.2. TDI A PPO2000 endcapped with 2,4-TDI and extended with HMDA presented a room-temperature modulus of 4.4 MPa and an HS melting temperature of 103 ◦ C [50]. Moreover, polyurethane–urea materials based on PPO2300 with 2,4-TDI and various aliphatic diamines (C3, C6 and C12) were found to have SS Tg s at −54 ◦ C, crystalline flow temperatures at approximately 90 ◦ C, G -moduli at room temperature around 5 MPa and CS values at about 18% [55]. However, when employing 2,6-TDI, the flow temperatures with the diamines C3, C6 and C12 increased to 243, 269, 272 ◦ C, respectively. The SS Tg was −56 ◦ C, the G -moduli increased to 10 MPa and, at the same time, the elasticity was improved (CS values of 14, 7 and 8%). The fact that the 2,6-TDI is less irregular than its 2,4 counterpart had a direct effect on the physical and mechanical properties. When

employing a TDI-endcapped PTMO, extended by aromatic diamines with an irregular structure, copolymers with an HS Tg in the range of 200–250 ◦ C were obtained [51]. 2.3.3. Aliphatic diisocyanates Fully aliphatic polyurethane–ureas have been studied for possible biomedical applications [33,34]. A poly(␧caprolactone) with a molecular weight of 2000 g/mol was used for the SS. The urethane–urea segments comprised aliphatic DII (l-lysine DI, BDI, HDI) with butane diamine as the extender and were monodisperse in length. The low temperature properties of these materials were comparable. The diurethane–diurea segments were crystalline, and displayed sharp melting transitions. The HS melting temperature was 91 ◦ C for the lysine DI, 129 ◦ C for HDI and 218 ◦ C for BDI. These are very significant differences and it is unclear why BDI demonstrated such a high HS melting temperature as opposed to the HDI-based copolymers. The storage E-modulus at 50 ◦ C (i.e., above the PCL melting temperature) was quite high for the HDI and BDI copolymers, i.e., 55 and 100 MPa, respectively, and the moduli were little dependent on temperature. Moreover, under ambient conditions, the Young moduli for these copolymers were comparable, i.e., approximately 40 MPa. The TS values at 50% strain were on the order of 20–30% of the applied strain. 2.4. Urethanes with other extenders Starting from DII-endcapped prepolymers, also other diol and diamine extenders have been used (Fig. 13). Investigations have been performed on DIIendcapped prepolymers extended with OH-R2 -OH units. For the latter, diol–diester and diol–diamide units were employed, enabling the formation of tetraurethane–diester and tetra-urethane–diamide segments, respectively [56]. Also NH2 -R3 -NH2 diamine–diamides extenders have been utilized, thus giving rise to diurethane–diurea–diamide segments [56–59]. Various properties for PTMO1270 –HDI–extender–HDI copolymers are presented in Table 4 [56].

Fig. 13. Extenders for isocyanate-endcapped prepolymers.

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Table 4 Polyurethanes constituted of PTMO1270 –HDI endcapped prepolymers and various types of extenders [56]. Extendera

inh (dl/g)

HS (wt%)

NH groups

Tg (◦ C)

 G25 ◦ C (MPa)

Tflow (◦ C)

CS (%)

Urethane

Urea

Urethane

1.8 1.2 1.1

4 4 4

0 0 0

0 0 0

−65 −66 −67

39 31 23

100 95 80

26 26 30

32 39 33

1.2 1.1 1.0

4 4 4

0 0 0

2 2 2

−67 −65 −66

44 56 34

155 160 140

42 37 26

32 33 35 40

1.6 1.2 1.4 1.3

2 2 2 2

2 2 2 2

2 2 2 2

−68 −67 −67 −65

22 49 49 47

175 245 225 195

18 21 25 14

Diol–diester –u 3e Te 3u – –u 4e Te 4u – –u 6e Te 6u –

33 34 36

Diol–diamide –u 2a Ta 2u – –u 3a Ta 3u – –u 3a Aa 3u – Diamine–diamide –ur 3a Aa 3ur – –ur 4a Aa 4ur – –ur 6a Aa 6ur – –ur 12a Aa 12ur –

a The figures correspond to the number of methylene groups; T stands for terephthalic and A for adipic; the subscripts u, ur, e and a correspond to urethane, urea, ester and amide, respectively.

The extenders were prepared prior to the copolymer synthesis. By using these longer extenders, the HS content was increased. The tetra-urethane–diester HS had a melting temperature of 80–90 ◦ C, a G storage modulus of about 30 MPa (E-modulus ∼ 90 MPa) and a CS of about 28%. Consequently, by incorporating the ester units in the extender, the modulus had increased as compared to BDO (Table 2), but not the melting temperature. The use of a diol–diamide (3T3) led to the formation of a tetra-urethane–diamide segment. The melting temperature now increased to 150 ◦ C and, at the same time, the G storage modulus increased to approximately 50 MPa (E-modulus ∼ 150 MPa), and also the CS values were higher, about 37%.

2.4.1. Diamine–diamide Diamine–diamide units have also been used as extenders (Table 4). These were based on either terephthalic acid (T) or adipic acid (A) with several aliphatic diamines, i.e., C3 , C4 , C6 and C12 [56]. The diamine–diamide units are called xTx and xAx wherein x stands for the number of methylene units in the diamine (Table 4). The amine groups with the isocyanate moieties formed urea units leading to diurethene–diurea–diamide segments being obtained. The extenders with even numbers of diamines in the diamine–diamides had G moduli of 49 MPa (E-modulus ∼145 MPa) at room temperature and HS melting temperatures that decreased with an increasing diamine methylene length from 245 ◦ C (4T4), 225 ◦ C (6A6), to 195 ◦ C (12A12). The CS value was found to be the lowest for 12A12 (14%). The 3A3 extender gave rise to a lower G -modulus, i.e.,

22 MPa, and a lower melting temperature, i.e., 175 ◦ C, while the CS values remained unchanged. Studies have also been carried out on the effect of the length x in –6-(T-6)x -diamine–amide extender, by varying x from zero to two (Table 5) [55]. By increasing x, in the –6-(T-6)x –, the HS melting temperature and the modulus were increased and the elasticity was improved. All of these copolymers displayed rubber plateaus that were independent of temperature. Also the melt stability at 200 ◦ C of the copolymer with the diamine–diamide extender seemed to be superior to that of a commercial polyurethane [55]. The type of DII coupling unit was also varied (TDI, MDI, HDI) with a 6A6 diamine–diamide extender (Table 5) [57]. In all these systems, the amide crystallinity was high, e.g., 70–80%, and the Tg of the PTMO phase was very low (∼−65 ◦ C). The room-temperature shear modulus (G ) value for TDI and MDI was found to be 16 MPa, whereas it was 32 MPa for HDI. The melting temperatures of the HS increased in the following order: TDI (165 ◦ C), MDI (205 ◦ C) and HDI (225 ◦ C). The CS values were low, on the order of 12–18%, and with the 6A6 diamine–diamide extender, the difference between TDI and MDI was limited to the melt temperature. Surprisingly, the HDI copolymers displayed higher moduli and higher HS melt temperatures than their TDI and MDI counterparts, thus reflecting the effect of structural regularity.

2.4.2. LCP extenders Mesogenic units presenting a liquid crystalline (LCP) behavior have been used as extenders in polyurethanes

Table 5 Copolymers with various diurethane–diurea–amide HS. Polymer

x

HS (%)

inh (dl/g)

Tg (◦ C)

 G25 ◦ C (MPa)

Tflow (◦ C)

CS (%)

Ref.

PPO2300 –TDI2,4 –6-TDI2,4 PPO2300 –TDI2,4 –6T6–TDI2,4 PPO2300 –TDI2,4 –6T6T6–TD2,4

0 1 2

19 26.7 32.9

0.58 0.6 0.35

−51 −53 −55

4 13 16

103 169 264

15 10 6

[55] [55] [55]

PTMO1886 –TDI2,4 –6A6–TDI2,4 PTMO2209 –MDI–6A6–MDI PTMO2000 –HDI–6A6–HDI

1 1 1

26.8 27.6 25.3

1.4 1.3 1.6

−72 −63 −65

16 16 32

165 205 225

12 18 18

[57] [57] [57]

R.J. Gaymans / Progress in Polymer Science 36 (2011) 713–748

Fig. 14. LCP extender of a TDI–PTMO–TDI prepolymer.

[60–63]. PU HS were prepared with a 50/50 mixture of TDI2,4/2,6 and a mesogenic bis(6-hydroxyhexoxy) biphenyl unit, forming liquid crystalline HS (Fig. 14). The HS was only one repeat unit long, and the LCP HS were found to be partly crystalline showing two transitions; one for the melting-mesophase and one for the mesophase-isotropic phase transition. With a decreasing LCP–HS content in the copolymer, the heat of isotropization of the meso phase, corrected for the LCP–HS content, was lowered. The PTMO SS apparently interfered with the LCP phase by mixing. The LCP phase was both thermodynamically and kinetically controlled. With the LCP extenders, the polyurethanes showed a gel point at the mesophaseisotropic phase transition [61]. 2.5. Amide Amide segments are more thermally stable, which allows higher copolymer melting temperatures without degradation during melt synthesis and processing. Amide groups can H-bond and also present a high polarity. The large difference in polarity between the amides and the SS signifies that even short amide segment lengths (i.e., a length of more than three repeat units, six amide groups) lead to liquid–liquid demixing [14,64]. Despite the multi-phase morphologies of the commercial polyamide segmented block copolymers, these materials have particularly good dynamic properties [64,65]. If the structural regularity of the amide unit is high, the melting temperature is elevated, as is the rate of crystallization. A fast

725

crystallization allows short cycle times during injection molding, and no after-cure is necessary. If the SS has amine end-groups, then with an acid or ester, an amide coupling is obtained between the SS and the HS. However, SS most often have hydroxylic endgroups, and the coupling with the acid functionalized HS thus occurs by means of an ester group. 2.5.1. Tetra-amide Studies have been performed on copolymers with tetra-amide HS based on amine-terminated PTMO750 and the diamide–dimethyl ester units (TXT-dimethyl) [66] (Fig. 15). The T stands for terephthalic, and the X represents the number of methylene units –(CH2 )x –. The methylene length, indicated by x in Fig. 15, was varied from 4 to 12 and only even-numbered methylene units were studied. As the number of amide units in the HS segment length was only four, liquid–liquid demixing was not observed. According to DMA, the copolymers presented only two transitions: a low Tg for the PTMO750 phase at −63 ◦ C and a melting temperature of the amide segment. The Tm increased with a decreasing methylene length, from 230 ◦ C for –(CH2 )12 – to 365 ◦ C for –(CH2 )4 –. For all these materials, the storage modulus G at room temperature was about 60 MPa (E-modulus ∼180 MPa). This is a high value, which is due to a relatively elevated HS concentration in the copolymers as a result of using the short PTMO750 length (32–38 wt%). The plateau modulus was temperature-dependent, thus suggesting that the amide segments had a narrow length distribution. 2.5.2. Diester-amide Copolymers with diester–diamide HS have been prepared starting from hydroxyl terminated polyethers and diester–diamide units. As aliphatic esters are susceptible

Fig. 15. Synthesis of copolymers with tetra-amide segments from amine-functionalized SS.

Fig. 16. Synthesis of copolymers with ester–amide segments.

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R.J. Gaymans / Progress in Polymer Science 36 (2011) 713–748

Table 6 Influence of the diester–amide segment length in –(PPO2300 –T-(6-T)x –)n – copolymers [70]. Polymer

HS (%)

No amide units

inh (dl/g)

Tg (◦ C)

 G25 ◦ C (MPa)

Tflow (◦ C)

CS (%)

PPO2300 –T6T PPO2300 –T6T6T PPO2300 –T6T6T6T

12 19 26

2 4 6

1.8 1.0 0.67

−57 −61 −61

4 28 30

85 194 264

29 13 11

Table 7 Some properties of PTMO1000 –diester–diamide as studied by DMA.

T2T T4T T6T TT BTB a b

HS (wt%)

inh (dl/g)

25 28 31 28 28

1.67 1.18 – 1.57 –

Tg (◦ C) −60 −65 – −69 −78a

 G25 ◦ C (MPa)

Tm (◦ C)

Ref.

25 65 – 44 35b

135 153 111 222 225a

[69] [68] [67] [71,72] [73–75]

From DSC. G corresponds to E /3.

to hydrolysis and thus most often not wanted in a polymer, esters based on aromatic moieties (e.g., terephthalate) are preferred, and the copolymerization is a polyesterification at a high temperature in the melt phase (Fig. 16) [67–84]. Also in these copolymers, the number of amide repeat units of the HS T(6T)x has been found to have a strong effect on the modulus and the HS melting temperature (Table 6). This is not only due to the length of the segment and the number of H-bonding units but partly also due to the increasing HS concentration. Melts of the polymers with diester–diamide and diester–tetra-amide segments were transparent during the polymerization, suggesting a homogeneous melt phase. The partial aromatic diester–diamide T6T had a low modulus, a low melting temperature and a high CS value. The diester–hexa-amide had a melting temperature of 264 ◦ C, too high for easy synthesis and melt processing. The polyether thermal stability is here the limiting factor. Furthermore, the copolymer with partial aromatic

diester–tetra-amide T6T6T segments presented good lowtemperature properties, a high melting temperature, a decent modulus and a low CS value. Other diester–diamide and diester–tetra-amide segments that have been studied are given in Fig. 17. The diester–diamide and diester–tetra-amide units were synthesized prior to the polymerization [67–84]. 2.5.3. Diester–diamide Several diester–diamide segments have been incorporated in copolymers, and the type of diamine in these segments has a strong effect on the melting transition. Various properties of PTMO1000 –diester–diamide copolymers are summarized in Table 7 [67–69,71–75]. The copolymers all presented elevated inherent viscosities and were thus of a high molecular weight and they could be processed from the melt without degradation. The Tg of the polyether phase for all these copolymers was low, which suggested that only small amounts of HS had been

Fig. 17. Various diester–diamide and diester–tetra-amide segments.

R.J. Gaymans / Progress in Polymer Science 36 (2011) 713–748

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Table 8 Various properties of PTMO1000 –diester–tetra-amides as studied by DMA.

T6T6T T8T8T T10T10T TPrTPrT T4A4T T6A6T

HS (wt%)

inh (dl/g)

Tg,SS (◦ C)

 G25 ◦ C (MPa)

Tm,HS (◦ C)

CS (%)

Ref.

35.0 37.2 39.2 32.5 31.9 34.2

2.2 1.6 2.4 2.0 2.2 2.5

−61 −62 −59 −67 −62 −60

87 72 60 33 70 51

235 215 210 220 215 185

18 20 15 28 19 14

77 82 82 83 82 81

dissolved in the PTMO1000 . The room-temperature storage modulus G was high, between 35 and 65 MPa. The main effect of the type of HS was on the HS Tm . The melting temperature increased from T6T to T4T; however, the T2T with the ethylene diamine was lower again. The Tm for both of the fully aromatic diamide units was moderately high. The elastic properties of the TT copolymers were interesting (CS 31%) and comparable to T6T in PPO–T6T (Table 6). 2.5.4. Diester–tetra-amide Several investigations have been performed on SBCPs with diester–tetra-amide segments, since they present a high melting temperature which is important for engineering applications. Some of these copolymers are given in Table 8 [70,77–83]. Also studied, although not presented in Table 8, were copolymers with the following tetra-amide segments: T2T2T [80], T3T3T [80], T7T7T [80] and TXpTXpT (where Xp = para-xyxlene) [78,79]. The glass transition temperatures of PTMO1000 were low in all cases, and the melting temperature of the tetra-amides increased with a decreasing diamine length. For the even diamines, it occurred at a higher temperature. When using an adipic group (A) with four methylene units, i.e., TxAxT, instead of a terephthalic group (T), i.e., TxTxT, the melting temperature of the segments was lowered, leading to the Tm for T6A6T being some 50 ◦ C below that of T6T6T (Table 8). The segments with the piperazine groups (–TPrTPrT–) were unable to form H-bonds. Nonetheless, due to a relatively rigid structure, they presented a Tm of 220 ◦ C. By using the p-xylene diamine (NH2 –CH2 –phenyl–CH2 –NH2 ), the melting temperature was 20 ◦ C higher than for the hexamethylene diamine in the similar TXTXT copolymer [78]. The TXTXT segments had a very regular structure and both the rate of crystallization and the crystallinities of the TXTXT segments were very high [77,80,83]. All H-bonding tetra-amides demonstrated high G moduli at room temperature (e.g., 51–70 MPa), whereas the corresponding values for the non-H-bonding TPrTPrT were lower. The compression set values for all copolymers were in the range of 14–20%, and somewhat higher for the non-H-bonding TPrTPrT. This was probably due to the crystallites of the latter, non H-bonded, apparently being more easily deformed. The diester–tetra-amide copolymers could be synthesized in the melt and processed from the melt while maintaining their monodisperse character [70,85]. As compared to the urethane and urea copolymers, the thermal stability of the copolymers with the amide segments was much higher.

2.5.5. Non-crystallizing amide segments If the amide segments were made with the irregular meta-xylene diamine (Xm ) and isophthalic acid (I) then the amide segments in the copolymer were amorphous [14]. The amide segments, –(Xm I)x –, phase separated at a length of x of three and longer and formed both nano and micro size particles. Thus, the (–(Xm I)x –PPO–)n materials had a complex morphology. The flow temperatures, the amide Tg s, increased with amide segment length, for x = 3 at 100 ◦ C to x = 10 at 195 ◦ C. The modulus of these materials at room temperature increased far less strong with amide concentration than with the crystalline amide segments. Surprisingly, the elastic properties as function of the modulus of these materials were as good as with the well-crystallizing amide segments [79]. 2.6. Type of the matrix segment A In the segmented block copolymers with monodisperse HS, the concentration of the SS is high. Moreover, the SS form the matrix phase, in which the HS crystallites are dispersed. This matrix phase can be fully amorphous or semi-crystalline and can either have a low or a high glass transition temperature (Fig. 2). Several matrix phase segments with a low Tg are given in Fig. 18. 2.6.1. Amorphous segments with a low Tg Typical amorphous segments of low Tg include PPO, poly(dimethyl siloxane) (PDMS), poly(ethylene-butylene) (a hydrogenated polybutadiene) (PEB), but also PTMO and PEO provided that they are of low molecular weight (Mn < 1000 g/mol). With such non-crystallizing SS, the copolymers have a two-phase morphology: an SS phase with a Tg at a low-temperature and an HS phase with a high Tm . However, PDMS and PEB are non-polar and the miscibility of these segments with H-bonding HS segments is generally very poor. A homogenous melt can only be obtained if both the SS and the HS are short. Short HS of monodisperse length are best suited for forming block copolymers with these non-polar SS. According to the literature, hydrophobic PEB3300 –TT copolymers have a Tg at −60 ◦ C, and a very low G -modulus at room temperature, i.e., 2 MPa [71]. These two-phase copolymers thus present excellent low temperature properties, good elastic behavior in combination with a wide service temperature range. However, their tensile properties are generally low since no strain crystallization takes place. On the other hand, strain crystallization may just occur with the short PTMO and PEO segments which are amorphous in the unoriented state but crystallize upon straining [86].

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length [70,78,80,84]. However, long PTMO or PEO segments (Mn > 2000 g/mol) have a melting temperature that is higher than room temperature, which is unfavorable for the room-temperature elastic properties. PEO copolymers are hydrophilic materials, and copolymers with PEO segments combined with monodisperse HS have a low contact angle with water [89], fast water vapor transport [90] and a high CO2 transport [91]. As a result of almost completely crystallized MD-HS, the PEO phase did not contain HS and this has led to an improvement of these properties in comparison to commercial materials.

Fig. 18. SS segments with a low Tg . The Tg and Tm values given are for segments with a molecular weight of about 2000 g/mol in a copolymer. The “am” means amorphous.

2.6.2. Semi-crystalline segments of low Tg Semi-crystalline segments of low Tg and Tm include PTMO (Mn > 1000 g/mol), PEO (Mn > 1000 g/mol), poly(butylene adipate) (PBA) and polycaprolactone. These segments crystallize upon cooling and can further crystallize upon straining. Such a strain-induced crystallization of the SS stabilizes the structure and results in the copolymers having higher fracture strains and stresses. Moreover, the SS melting temperature increases upon straining; an increase of 40 ◦ C has been observed upon straining PTMO [86]. The regular aliphatic esters PBA and polycaprolactone are easy crystallizable and have a Tm above room temperature, and their corresponding copolymers show decent mechanical properties [17,32–34,44,74,75]. However, the elastic properties with these crystallizable SS are only acceptable above the SS melting temperature (>50 ◦ C). Below their SS Tm , these copolymers demonstrate a shapememory behavior [87,88]. It would be nice if, at low strains, the SS were amorphous with good low-temperature properties and a high elasticity, combined with a strain-hardening effect at high strains to obtain decent tensile properties. Copolymers of PTMO1000–2000 and PEO1000–1500 demonstrate this behavior. The type of SS (PTMO/PPO/PEO) had a surprisingly small effect on the room-temperature modulus [84]. The elasticity of the copolymers improved with increasing SS

2.6.3. Extended soft segments With increasing length of the PTMO and PEO segments (Mn > 1000 g/mol) the melting temperature and crystallinity of those segments increase and thereby reduce the low temperature properties. Certainly a SS melting temperature above room temperature is not wanted. The SS crystallization of longer PTMO and PEO segments (e.g., Mn > 2000 g/mol) can be suppressed by copolymerization with another ether segment [26,92,93] or with the incorporation of structure-disturbing groups such as terephthalic [78,81,94,95], isophthalic [77] or DII moieties [58,59]. Long ether segments containing structure-disturbing groups can be obtained by reacting short ether segments with terephthalic or diisocyanate groups to long SS units (Fig. 18). This way, the SS lengths of PTMO1000 and PEO600 have been increased to 9.000 and even 20.000 g/mol, while the melting temperature was still below room temperature [77,81,94,95]. Consequently, the MDI-extended prepolymer MDI–(SS–MDI–)q remained a liquid at room temperature and the corresponding copolymers extended with SS have demonstrated a Tg only a few degrees higher than that of the pure SS [58]. PTMO and PEO extended with terephthalic groups also presented an SS Tg that was only 5 ◦ C higher. Thus, by incorporating disturbing groups in the crystallizable SS, long chain SS can be obtained without the SS melting temperature being too high. 2.6.4. Amorphous matrix segments with a high Tg Semi-crystalline homopolymers generally have a Tg /Tm ratio (in K) near 0.66. If the Tg of the homopolymer is higher than 150 ◦ C, then according to this 2/3 rule, the Tm will be well above 300 ◦ C and the crystallization window is about 200 ◦ C. However, for most polymers, Tm > 300 ◦ C is too high for a degradation-free melt synthesis and melt processing. It would thus be useful to have polymers with Tg above 150 ◦ C, but with a melting temperature below 300 ◦ C. The crystallization window for such polymers is then small and only very fast-crystallizing units would be able to crystallize upon cooling from the melt. As mentioned above, monodisperse tetra-amide segments are thermally stable, high melting and rapidly crystallizing units. It was thus of interest to study whether such segments, incorporated in a system with a high Tg , would maintain their ability to crystallize, thereby increasing the dimensional stability and the solvent resistance of the polymer. Studies have been carried out on copolymers of PPE, PC and PSU with the bisester tetra-amide units T6T6T (Fig. 19) [96–101]. Copolymers with T6T6T segments presented a T6T6T melting temperature that was not too high (>270 ◦ C).

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729

pared to the neat polyester the obtained polyester–T6T6T copolymer could be drawn at a higher temperature and to a higher draw ratio, thereby increasing the strength and strain of the drawn material. 2.7. Other copolymers

Fig. 19. Various segments with high Tg s.

Moreover, since the Tg s of the amorphous segments were high, the crystallization window for the T6T6T segments was small (<100 ◦ C) but the T6T6T segments could nonetheless crystallize from the melt upon cooling. The copolymers demonstrated an improved dimensional stability and solvent resistance, and surprisingly this even at T6T6T concentrations of less than 10 wt% [97,100,101]. 2.6.5. Semi-crystalline matrix segments with a high Tg and high Tm Segmented block copolymers with semi-crystalline matrix segments (such as PTMO) can strain harden, and high fracture stresses and strains are obtained (see Section 2.6.2). Investigations of whether the drawing behavior of a similar semi-crystalline polymer, but one with a high Tm could also be enhanced by modification with high melting monodisperse segments have been performed. It was found that the modification of a high melting polyester (PBT or PET) with a low melting diamide (TXT) did not form a system with two melting temperatures, but enhanced the polyester melting temperature, the crystallization rate and crystallinity of the polyester [102–104]. A higher melting temperature for the tetra-amide (TXTXT) segments and the incorporation of these TXTXT segments in a polyester were also studied. In a first evaluation of this principle, a polyester, poly(hexamethylene terephthalate) was modified with 10 wt% T6T6T segments (Fig. 2d) [105]. These copolymers demonstrated two melting temperatures: one for the polyester and a higher one for the T6T6T segments. Com-

Other copolymers studied include di-block (A–B) [72] and tri-block (A–B–A and B–A–B) [8,72,106,107] with monodisperse HS. Short monodisperse HS are effective in forming physical network structures at low concentrations. At low strains, the solid state properties of the tri- and the multi-block copolymer of PTMO–TT were comparable. The fracture properties decreased with a decreasing molecular weight of the copolymer like the multi-block copolymers [72]. The melt rheological behavior of the tri-block copolymers was interesting since the (amide) end-segments easily formed a gel [108–110]. Surprisingly, the di-block copolymer AB and the tri-block copolymer ABA with only one HS in the chain remained solid and presented reasonable properties [72,106]. In the context of another study, a segmented block copolymer was blended with unreactive HS [44]. The unreactive HS cocrystallized with the HS in the copolymer and thereby increased the modulus (due to the increased HS concentration) while the fracture properties of the copolymer only slowly decreased with increasing HS. Moreover, the HS in the copolymer could be modified with (bio)functional units attached to the HS, while maintaining decent mechanical properties (Fig. 20) [34]. Based on this principle, it is envisaged that biomaterials with specific functionalities can be readily fabricated. The functional units may be added either as part of low molecular weight compounds [34] or as mono functional end-units on the polymer chain. Also, as the MS HS ribbons have a high aspect ratio (Fig. 21) a percolating structure is easily obtained. If functional units are attached to the ribbons, then also the functional units form a percolation path at a low concentration. Membranes with a percolating structure of units groups can possibly form a molecule specific percolating path for transport, as occurs with biomembranes. These membranes might be useful for extracting one compound specifically from a mixture of compounds, and this particularly at low concentrations of that compound.

Fig. 20. A schematic representation of a diurea copolymer with functional end-groups (gray spheres). Reproduced with permission from ref. [34], copyright (2006) American Chemical Society.

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Mn =

 NM i i

(2)

Ni

The MD-HS units in the copolymers are usually short in length (Mn < 1000 g/mol, x < 3 repeat units). Units such as T(XT)x amide units are synthesized prior to the polymerization from two difunctional monomers with an excess of one of the components. In T(XT)x the T stands for terephthalic and X for the number of methylene units in the diamine part. The stoichiometry of the T and X compounds (i.e., the X/T ratio) is thus smaller than unity (r < 1). The polydispersity of the segments with excess of one component may be calculated with the most probable distribution function of Flory [111], assuming that the reactivities of the endgroups of the difunctional compounds are equal, Eqs. (3) and (4). wn = nr ((n−1)/2) Nn =

wn n

(1 − r)2 (1 + r)

(3) (4)

Here,

Fig. 21. A cartoon of the membrane (grey area) with nano-ribbons with a high aspect ratio (in green) with functional units (in red), forming a nano percolating structure of the functional groups (red dots) with a group specific transport of the (triangle sign) molecules. The transport is from top to bottom. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)

3. Properties 3.1. Introduction For SBCPs with H-bonding units to have with SS a homogeneous melt phase the HS segments should be short in length (Mn,HS < 1000 g/mol), as otherwise liquid–liquid demixing will usually take place. Short, monodisperse HS can exhibit effective crystallization. The HS length distribution is important not only for obtaining a homogeneous melt, but also for fast crystallization, high crystallinity of the HS and high aspect ratio of the crystallites. The polydispersity of the hard segments has a strong effect on the morphology and properties of the copolymers. 3.2. Polydispersity Segmented copolymers are mostly synthesized by a step-growth mechanism and such step-growth polymers ideally have a molecular weight distribution for which Mw /Mn = 2. Most linear polyesters, amides and urethanes present this distribution, which is also called the polydispersity index (PDI) [111]. The number average (Mn ) and weight average (Mw ) of the hard segments can be calculated according to the following well-known formulae (Eqs. (1) and (2)): Mw =

 N M2  i i Ni Mi

(1)

wn = weight fraction (%/100) of n-mer with n = odd numbers (1, 3, 5, etc.), n = n-mer (n is odd numbered, referring to the odd specimen), r = stoichiometric balance (<1), Nn = number average (%/100) of n-mer with n = odd numbers (1, 3, 5, etc.). Using these equations, a graph can be constructed of the concentration of the segments with a particular length n. From these data the polydispersity can be calculated. According to this relationship, the polydispersity decreased with an increasing excess and thus a decreasing degree of polymerization. However, this effect is not very strong. For the random polymer with a repeat length of two, the stoichiometric ratio r is 2/3. The calculated weight distribution of these “randomly” synthesized hard segments is 1.89 [85,112]. In the above equations, equal reactivity between the first and second reactive group was assumed. However, the reactivity of the second reactive group for diisocyanates and diacids is often lower and this difference might even be large. For the MDI, the reactivity difference K was a factor 2.9, whereas it was 11.9 for the 2,4-TDI [12,22,23]. If the reactivity of the first group k1 is higher than that of the second group k2 , then a majority of the first group has reacted before the second group takes part in the reaction. In stoichiometrically balanced systems (r = 1) with an unequal reactivity (K = k1 /k2 > 1), the polydispersity hardly changes. However, if in addition to an uneven reactivity (K > 1) also the stoichiometric ratio r is smaller than one, the polydispersity is reduced further [23]. Another effect in SBCPs involves longer H-bonding segments becoming incompatible with the matrix phase (or solvent) and separating. This phase separation of the longer HS segments limits further reaction of these segments and thereby reduces the PDI. For instance, if one assumes that

R.J. Gaymans / Progress in Polymer Science 36 (2011) 713–748

segments longer than 6 repeat units are not formed for r = 2/3 and K = 1, then the PDI changed from 1.89 to 1.65. The combination of one compound in excess, a high K value and phase separation at long length was studied in the synthesis of T(6T)2 -dimethyl from DMT and HMDA at a ratio of 3:2 in an NMP solution [85,112]. It was observed that the PDI for T(6T)2 -dimethyl was only 1.24, which is significantly lower than the PDI value of 2 that is normally obtained for a high molecular weight polymer. In fact, the polydispersity in the one-pot synthesis of short segments can thus be rather small and depends on whether the longer units crystallize during the synthesis. There are several studies where copolymers with monodisperse and polydisperse HS have been compared [2,3,6,8–9,12,14,70,78,85]. The effect of the HS molecular weight distribution on the copolymer properties has been examined in detail for amide segments [85]. SBCPs with a varying HS distribution were prepared by using a mixture of several MD starting blocks (T6T, T6A6T and T6A6A6T) and PTMO2000 segments and in this way the PDI was varied from 1.00 to 1.09. A copolymer with the same overall composition, but with randomly prepared HS, demonstrated a PDI of 1.2. The distribution of the “monodisperse” starting units can be studied with NMR [71,76], GLC [8,9] and MALDI-ToF [85,112], but the determination of the distribution of segments in a copolymer is extremely difficult. Nevertheless, if the HS can be isolated by hydrolysis of the linkage between the HS and the SS of the copolymer, it becomes possible to determine the molecular weight distribution of the isolated HS [85,112]. For copolymers with diester–amide HS, such as PPO–T6T6T and PTMO–T6A6T copolymers, the ester linkage was hydrolyzed and the HS isolated. It was found that the HS in the copolymer had the same distribution as the starting units. The T6T6T and T6A6T HS distribution of the amide segments had not increased upon performing a high-temperature melt polymerization. The distributions of the initial amide units were thus representative of the T6T6T and T6A6T segment distribution in the copolymers, even after several hours at 250 ◦ C [85,112]. The rate of trans-reactions of amide groups is small and the thermal stability high. Apparently, the HS length distribution of short segments with unequal reactivity of the reacting groups is much smaller than 2 and is thus referred to here as “polydisperse”. The term “random distribution” would suggest a distribution of 2. On the other hand, if the segment length distribution is less than 1.03 it is accepted as “monodisperse” (MD), and if the distribution is expected to be better than polydisperse, but not monodisperse, the term “narrowly dispersed” is employed. 3.3. Melting temperature of HS (DSC and DMA) The melting transitions can be studied by DSC and DMA: the flow temperatures (Tflow ) as measured by DMA correspond well with the melting temperatures as measured by DSC [26,68,70–72,77,78,81,84]. For SBCPs with low amounts of crystallized HS, the value of Tm , is often difficult to determine by DSC. In such cases, the information from DMA is particular relevant. Even at HS concentrations

731

Fig. 22. A DSC trace of PTMO2016 –Pyp–(BDO–Pyp)6 –polyurethane with polydisperse hard segments: top line, first heating; bottom line second heating after annealing. Reproduced with permission from ref. [9], copyright (1989) Springer–Plenum.

as low as 2–6 wt%, a flow transition has been observed with DMA [26,81,94]. The melting temperatures of compounds depend on the ratio of the melting enthalpy (Hm ) and the melting entropy (Sm ) (Eq. (5)). Tm =

Hm Sm

(5)

For polyurethanes, polyureas and polyamides, the high melting temperatures of the H-bonded compounds are not high due to an elevated melting enthalpy (Hm ), but due to the low melting entropies [113,114]. Apparently, the chain flexibility of the H-bonded units does not change much upon melting as they remain H-bonded even in the melt. A copolymer with a polydisperse distribution of the HS length showed multiple melting transitions, giving rise to a broad DSC melting transition (Fig. 22) [8,9]. Upon annealing this copolymer, the peaks sharpened and each peak could be linked to a particular segment length [8,9]. The copolymers with monodisperse HS demonstrated sharper melting transitions [8,9,85]. However, when the repeat length was longer than four units, the monodisperse units no longer stacked neatly, giving rise to conformational disorder (Fig. 23). Thus, for long repeat units of monodisperse length (n > 4), due to conformational disorder, the stacking of the chains becomes less regular. As a result of this the crystals are no longer uniform in thickness and the advantage of having MD-HS has disappeared [9]. The melting temperatures of the model compounds and their copolymers generally increase with the number of repeat units [2,3,8,9,12,24,43,115]. The effects of the HS length and concentration (soft segment length) are clearly illustrated in Fig. 24. This effect of the HS chain length on the melting has, for model compounds, been described by the Flory–Vrij theory [8,9], which takes into account the surface effects of the crystals. For polymers, an often utilized relationship between the crystallite thickness and the melting temper-

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Fig. 23. Segmented block copolymers with monodisperse hard segments. For a repeat length x > 3, conformational disordered stacking takes place. Reproduced with permission from ref. [9], copyright (1989) Springer–Plenum.

ature of polymers with a lamellar morphology is that of Hoffman–Weeks (Eq. (6)) [25].



∞ Tm = Tm 1−

1 Hm

 2  e

l

(6)

PTMO segments, and that of the PPO-segments was somewhere in between. For a semi-crystalline polymer in presence of a solvent or plasticizer and for copolymers, the melting point depression has been described according to the following relation by Flory (Eq. (7)) [116].

The melting temperature Tm depends on the ideal melt∞ of a fully extended chain, the crystallite ing temperature Tm thickness l, the surface tension  e and the heat of melting Hm . For SBCPs with short segments, these segments are in the crystals present in their extended from and the length of the segment is the crystallite thickness l. Increasing the HS concentration (decreasing the SS concentration) also increased the melting temperature (Fig. 24). The HS melting temperature has also been found to depend on the HS–SS interaction [31,70,71,77,84]. Increasing the length of the flexible end-groups of the model polyurethanes decreased in the melting temperature [8,9,12,24,43,115]. Also, in SBCPs with MD-HS, increasing the SS length (decreasing HS concentration) has given rise to a decrease of the melting temperature (flow temperature) (Fig. 25) [31,70,71,75,77,84]. The melting temperature in the copolymers decreased when decreasing the amide concentration (increasing ether length) and varying the type of polyether. The amide melting temperature changed stronger with PEO than with

0 where Tm is the melting temperature of the copolymer, Tm is the melting temperature of the crystalline homopolymer, R is the gas constant, Hm is the melting enthalpy of the crystalline segment, Vp is the molar volume of the crystalline segment, Vs is the molar volume of the solvent,  is the Flory interaction parameter and vs is the volume fraction of the solvent. According to this equation, the depression of the copolymer’s melting temperature due to the presence of a solvent depends on the melting enthalpy, the ratio of molar volumes and/or the Flory interaction parameter. Although the Flory relationship was proposed for low contents of solvent, it has been applied to SBCPs with high contents of SS. When employing the equation for segmented block copolymers, two problems are encountered: the SS content is mostly high (>50 vol%) and it is difficult to define the molar volume of the SS (Vs ) and HS (Vp ) as these

Fig. 24. The HS melting temperature as a function of the amide concentration of –(PPO–T(6T)x )n – for various HS lengths x: 䊉, 1; , 2; , 3 [data from 70].

Fig. 25. The hard segment melting temperature as a function of the ether concentration of polyether–T6T6T with varying polyether segments: , PTMO; , PPO;  PEO (data from respectively 70, 77, 84).

 Vp  1 1 R − 0 = × × vs − v2s Tm V H Tm s m

(7)

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733

Table 9 The thermal transition properties of the amide segment for PTMO2000 –T6A6T (21 wt% HS) with a varying HS polydispersity [85]. Polydispersity

DMA

DSC ◦

1.0 1.01 1.03 1.04 1.07 1.09 1.2 a b

FTIR ◦

◦

◦

Tflow ( C)

Tm2 ( C)

Tc1 ( C)

Tm2 − Tc1 ( C)

Hm2 (J/g T6A6T)

Xc (%)

Xc (%)

180 180 175 180 185 185 165

183 180 174 182 190 184 –a

164 163 154 159 165 154 –a

19 17 20 23 25 30 –

27 26 24 22 19 7b –

93 89 83 75 65 24b –

92 84 85 74 80 80 71

Not measurable. Broad DSC traces, difficult to integrate.

segments have completely different chemical structures and repeat lengths. The effect of polydispersity (1.00–1.09) on the main melting in the PTMO2000 –T6A6T temperature as measured by DSC and DMA was found to be small (Table 9). However, no melting peak was observed with DSC for the copolymer with polydisperse hard segments (PDI = 1.2) and the Tflow obtained by DMA was 20 ◦ C lower. Certain studies have nevertheless indicated that the melting of the monodisperse and polydisperse HS is similar [30]. 3.4. Crystallinity (DSC and FTIR) PBT-based SBCPs with PTMO segments with a length of 1000 g/mol and 63 wt% PBT segments had a PBT crystallinity of about 33%, almost as high as for the neat PBT [117]. The SS (PTMO) apparently had little effect on the PBT crystallinity. The crystallinity decreased with a PBTsegment length shorter than 15 repeat units. The HS crystallinity of MDI-based polyurethanes prepared with a polydisperse HS distribution is often low, e.g., around 20%, similar to low crystallinity obtained with monodisperse MDI HS [4,30]. The crystallinity was much higher for the more regular HMDI- or p-PDIbased PUs [28,30]. Polyurethanes with short monodisperse piperazine–BDO segments also have a very high crystallinity (75%) [9]. The crystallinity of the HS depends on the structural regularity (symmetry) of the HS chain, the crystallization window and the melt viscosity, but the H-bonding has very little effect [9,83]. Similarly, the crystallinity of linear polyethylene is higher than that of the nylons. Normally, wide angle X-ray scattering (WAXS) is a good method to determine the crystallinity of polymers. However, hardly any crystalline scattering takes place for SBCPs with a low concentration of the HS and a short length of the HS chains. Neither is it possible to obtain the crystallinities for SBCPs from density measurements. Thus, the crystallinity of SBCPs is usually studied by DSC. The MD-HS crystallinities were often very high (>70%) and this also for the non-H-bonded HS [9,83]. Unfortunately, if the HS contents are low (e.g., less than 15 wt%), it is often impossible to observe the melting endotherm [26,71,77,81]. The crystallinities of HS in SBCPs may be determined by FTIR, provided that specific bands sensitive to the packing are present. In this respect, the position of the carbonyl band is a very good indicator, as this band is sensitive to

the H-bonding distance [20,26,28,58,59,80,81,93,102,115]. The shift in the carbonyl band has been studied in great detail on MDI polyurethanes; however, due to the polymorphic structure of these polyurethanes it is difficult to determine the crystallinity in this way. For semiaromatic amides, the carbonyl band in the crystalline state is 1625 cm−1 as opposed to 1670 cm−1 in the amorphous state. Since the distribution of H-bonding distances in the amorphous phase is wider, the 1670 cm−1 peak is broader. Upon melting a sample of PEO–T6T6T, the crystalline peak at 1625 cm−1 disappeared and the 1670 cm−1 peak enlarged (Fig. 26) [84]. The change in crystallinity upon heating and cooling could be determined from this shift [26,81,84]. The FTIR bands for the amide carbonyl groups are so specific that the crystallinities of amide HS in SBCPs could be properly measured for which the HS concentration was as low as 4 wt% [72,81,84]. With the TT, T6T6T and T6A6T HS, the crystallinity was very high (70–90%) and remained high even near the melting temperature (Fig. 27) [81]. Upon cooling, the crystallinity increased rapidly and the super-cooling effect was small. The amide crystallinity in amide-extended PU (e.g., –TDI–6T6–TDI–) was also high [55,57–59]. With increasing polydispersity the amide crystallinity of PTMO2000 –T6A6T, as determined by DSC and FTIR,

Fig. 26. FTIR spectra of PEO2000 –T6T6T in the 1580–1760 cm−1 wavelength region at various temperatures: , 50; , 100; , 150; , 170; 䊉, 190 ◦ C. Reproduced with permission from ref. [84], copyright (2007) Wiley & Sons.

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Fig. 27. Amide crystallinity of PTMO2000 –T6A6T as determined by FTIR in a heating-cooling cycle at 5 ◦ C/min. Reproduced with permission from ref. [81], copyright (2007) Wiley & Sons.

decreased (Table 9). However, the sample with a PDI of 1.2 had an FTIR-determined crystallinity still as high as 71%. The crystallinities as observed with FTIR correspond well with those obtained by DSC [81]. The HS crystallinity is normally not influenced by the type of SS, however for PEO an unusual effect has been observed. With PEO segments in a PEO2000 –T6T6T SBCP, the T6T6T heat of fusion decreased although the T6T6T crystallinity, according to FTIR, remained high [84]. These results suggest that, with PEO segments present, the heat of fusion was lowered due to a mixing exotherm as a result of a strong PEO–T6T6T interaction.

Fig. 28. The difference between the melting temperature and the onset of crystallization (Tm − Tc ) of the amide segments as a function of the polydispersity of T6A6T–PTMO2000 . The scan rate was 20 ◦ C/min. Reproduced with permission from ref. [85], copyright (2008) Wiley & Sons.

In principle, super-cooling values can also be obtained from thermo modulated FTIR (Fig. 27) [26,81,84] and SAXS [81]. Both these methods also confirmed the low supercooling values of the MD-HS measured by DSC. The reason for the very high rates of crystallization of copolymers comprised of short MD-HS seems to be the ease of ordering of these short segments. Another effect might be a pre-ordering of the H-bonded segments in the melt. However, according to SAXS studies copolymers in the melt, no liquid crystalline ordering was observed for T6T6T and T6A6T segments [26,81]. Furthermore, non-Hbonded MD amide segments were found to crystallize very fast too [83].

3.5. Rate of crystallization 3.6. Thermal stability The rate of crystallization is a very important property as it has a direct influence on the processability of polymers. A high rate of crystallization allows fast demolding of the materials, with little secondary crystallization. The properties of the molded materials thus change little over time. During melt-spinning of these elastic materials, a high rate of crystallization also permits higher spinning speeds. The rate of crystallization depends on the structural regularity, the window of crystallization and the melt viscosity [118]. The rate of crystallization is often studied by isothermal spherulite growth measurements. However, SBCPs with MD-HS with concentrations below 20 wt% do not have a spherulitic structure, and moreover are transparent materials [26,70,71,77,81]. Isothermal DSC measurements might be an option, but due to the often low concentrations of the HS, the crystallization peaks are usually too weak to give accurate results. In this case, dynamic DSC measurements are a possibility. These involve measuring the temperature difference between the melting peak and the onset of crystallization (Tm − Tc ) at a particular scan rate. SBCPs with MD-HS have been found to present very low super-cooling values; much lower than those of homopolymers [26,70,71,77,81]. The T6A6T MD-HS in the copolymer T6A6T–PTMO2000 demonstrated a super-cooling of only 18 ◦ C (Table 9, Fig. 28). With increasing polydispersity, the super-cooling values increased and consequently the rate of crystallization was lowered [85].

There are two aspects of the thermal stability of SBCPs with MD-HS: the chemical stability on heating in the presence or absence of oxygen, and trans-reactions. Upon chemical degradation, the molecular weight decreases and color formation takes place. This color formation occurs to a lesser degree with aliphatic urethanes, ureas and amides. The trans-reactions randomize the segment length distribution from monodisperse to polydisperse (Fig. 29A). For melts of polyurethanes and polyureas, i.e., above 190–200 ◦ C, these trans-reactions are fast [9]. Polyamides, on the other hand, have much slower trans-reactions and MD units are stable during melt synthesis at 250 ◦ C and melt processing [85,112]. A special case involves diurethanes, diurea, diurethane–diurea and diurethane–diurea–diamides (Fig. 29B). For these materials, any trans-reactions that might take place between two urethane (or urea) groups would not randomize a MD structure since urethane (or both urea) groups are identical. As a result, the HS with diurethane–diurea–diamide units have been found to be more stable than polyurethanes in general [55]. 3.7. Morphology The crystalline morphology of the HS has a strong effect on the mechanical behavior of the copolymers. A simpli-

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Fig. 29. Effect of trans-reactions on dispersity: A, tetra-urethane; B, diurethane.

fied SBCP has a two-phase structure: a low-Tg amorphous matrix phase of the SS and a dispersed crystalline phase of the HS. The structure of the HS crystallites has been extensively investigated by several methods, including AFM, TEM, SAXS and WAXS. When examining microtomed bulk PU samples, crystalline ribbons, threads or fibrils are often observed, dispersed in the soft matrix [119–121] (Fig. 11). A schematic representation of such a crystalline HS ribbon structure is given in Fig. 30. The crystallized ribbons have one long axis in the Hbonding direction (a) as well as two short axes of width (b) and thickness (c). The HS chain runs in the direction of the thickness (c), and with MD-HS, the ribbon thickness is expected to be the MD-HS extended length. The crystallites are thus expected to have a uniform thickness (Fig. 1). In the width direction (b) the chains are bonded with a van der Waals interaction. The fact that the a-axis is much longer than the b-axis must be due to the higher crystallization rate in the H-bonding direction. The short monodisperse segments show an extended chain length on the order of 1–5 nm and this is expected to be the thickness of the crystallites (c) [8,9]. 3.7.1. TEM In TEM analysis on thin annealed films from diluted solution on PUs sometimes lamellae and, in other cases, ribbons or threads of the crystalline HS are seen. On microtomed bulk PU samples, ribbons, threads or fibrils are often observed, dispersed in the soft matrix [119–121]. The fact that the crystallites in annealed films from solution more often have a lamellar structure suggests that the width of

Fig. 30. A schematic representation of the morphology of crystalline ribbons of a PU copolymer of HS with a narrow distribution. Reproduced with permission from ref. [119], copyright (1975) Francis & Taylor.

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Fig. 31. TEM analysis on an annealed cast thin film (50 nm) of PPO–T6T6T stained with OsO4 , of a polymer described in reference [70].

the crystallites (b) depends on the crystallization conditions. For SBCPs with MD-HS, very few TEM analyses have been performed on bulk samples. Most characterizations have been carried out on very thin cast films (50 nm) as for example in Fig. 31 [70,78]. For this PPO–T6T6T sample ribbons with a length (a) of 2–5 ␮m and a width (b) of about 100 nm were observed. The ribbon thickness (c) could not be determined, but as the extended length of the MD-HS (T6T6T) was 3.7 nm, the ribbon thickness was expected to be about the same. The relatively high width might have been due to the sample preparation method in which a very thin solution cast film was annealed. The observed long length of the ribbons was typical for MD-HS [26,42,72,81,122]. It could thus be concluded that the MD-HS crystallized very fast in the H-bonding direction. 3.7.2. AFM The structure of the crystallites in the copolymers has often been studied by AFM in the phase angle mode, and hard crystallites in a soft matrix near the surface are visualized by this method. The surface can be analyzed to a depth of about 100 nm, or less if the hard phase is encountered first. Generally, the surfaces are microtomed bulk samples or cast films, 10–100 ␮m thick, much thicker than those used for TEM analysis. When performing AFM on bulk samples and solution cast (non-heat-treated) films of copolymers with MD-HS, one generally observes invariable nano-ribbons [26,30,72,81,122], threads [42] or cylinders [9] (Fig. 11). These ribbons are branched to a very slight degree, extremely long (i.e., 300–1000 nm) and intermingled like in a dish of cooked spaghetti. Moreover, they present a high aspect ratio of >100. The structure of the ribbons or threads also suggests that the crystallites have two short and one large dimension, as displayed in Fig. 30. When lowering the concentration of the monodisperse HS, the ribbon structure did not seem to change. Long ribbons have been seen even at very low HS contents (3 wt%) (Fig. 32) [81].

Fig. 32. AFM of PTMO20000 –T6A6T with 3 wt % of HS. Reproduced with permission from ref. [81], copyright (2007) Wiley & Sons.

The crystals of the MD-HS have a ribbon-like morphology, whereby the width and the thickness are of the same size. This means for the diurea, diamide and tetra-amide HS that have extended chain lengths of about 1, 2 and 4 nm that their crystallite thickness (c) is of that order too. The small lateral dimensions of the ribbons (b) must be due to the relative slow growth rate in the van der Waals bonding direction. On annealing, the width increased as can also be seen in Fig. 33b, compared to 33a. Most often, the ribbons are linear, not branched, and fairly straight (Figs. 11, 32 and 33a). However, when the chemical structure of the units in the HS was less linear (as for MDI or TDI units), curved crystallites were observed (Fig. 34) [58,59]. Curved crystallites are expected to be less effective in reinforcing the matrix material. Crystalline ribbons have also been observed with nonH-bonding MD-HS, but these ribbons seem to be shorter in length [83]. 3.7.3. WAXS With short MD-HS, the number of repeat units in the HS is low, and as a result the scattering intensity in the c direction of the crystalline ribbons is very weak [12,26,86]. The size of the crystallites in the width direction (b) is also small, but the number of repeat units is larger in this direction. Moreover, in this direction in principle it is possible to obtain the size from the line broadening in WAXS. However, WAXS analysis on polyurethane SMBC with short MD-HS demonstrates very weak peaks for the MD-HS – often too weak to be analyzed [26]. 3.7.4. SAXS The structures of SBCPs with crystallized HS have been extensively analyzed by SAXS [12,26,37,47,81,123]. With SAXS it is possible to study the repeat distance at a

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Fig. 34. An AFM image of a PTMO3900 –TDI–6T6–TDI copolymer with curved ribbon crystallites. Reproduced with permission from ref. [59], copyright (2009) Wiley–VCH.

Fig. 33. The effect of annealing on the structure, as studied by AFM of a PTMO2000 –TxTxT; a, cast; b, cast and annealed.

2–100 nm scale. When analyzing the SAXS data of SBCPs, it is often assumed that the crystallites have a lamellar form, with one small dimension in the chain direction and two large dimensions. However, monodisperse HS crystallites rarely have a lamellar morphology, but rather it resembles that of a ribbon (Fig. 30). The ribbons have a long axis (a) of several hundred nanometers and this is outside the measuring range of this method. The ribbons have two short axis b and c of which the width (b) changes upon annealing (Fig. 33). Consequently, the long period in width (b) and the crystallite thickness (c) direction can be studied. A potential problem with crystalline ribbons is the expected overlapping of the long periods of the b and c

directions. Analyzing this 3D morphology with the 2D SAXS method is difficult as the SAXS diffraction gives an average of the long periods in the c and b directions. The long period in the c and b directions can be separately studied on oriented samples with the X-ray beam in the a direction. However, this analysis has yet to be reported for SBCPs with MDHS. Another complication is the fact that the crystallites of an SBCP are broken on drawing, leading to the sizes in the a and b directions not being what they were in the undrawn unoriented state [17,18,37,124,125]. The long period of a ribbon-like structure is thus an average of the values in the b and c directions, particularly as the crystalline ribbons are not neatly stacked but disordered (Figs. 11, 32, 33a and 34). The distances between crystallites are expected to increase with the square root of the SS length, in line with the endto-end distance of randomly coiled polymers [124]. SAXS analysis has also been performed on copolymers with diurea HS and PTMO1000 SS. The diurea segments have an extended length in the order of 1 nm. For BDI, a long spacing of 6.0 nm as well as an interdomain spacing of 4.5 nm has been reported [34]. The observed interdomain spacings for p-PDI, m-PDI, CHDI and HDI were 5.5, 5.5, 7.9 and 7.9 nm, respectively [36]. These are remarkable differences considering that the diurea unit length and SS length hardly changed. Interesting information can also be obtained from thermal modulated SAXS [26,81]. When raising the temperature, the long period of SBCPs with MD-HS increases only slightly up to near the melt temperature, after which the increase becomes rapid (Fig. 35). Upon cooling, the long period decreases very rapidly, causing the super-cooling effect to be small. In the melt, the coherent scattering disappeared, suggesting that no ordering was present [26,81].

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ture dependence of the modulus at the rubbery plateau, in the Tg − Tm region, is given by G , expressed in % change in G per ◦ C. Often, the SS has a regular structure (e.g., in PTMO) with an ability to crystallize, and these crystallites also reinforce the amorphous SS phase. After melting of the SS crystallites, the modulus reaches a stable rubber plateau. The SS melting temperature coincides with the onset of the rubber plateau which is denoted the flex temperature Tflex .

Fig. 35. The long period as measured by SAXS of PTMO2000 –T6A6T in a thermal cycle (Tm = 180 ◦ C). Reproduced with permission from ref. [81], copyright (2007) Wiley & Sons.

3.8. Dynamic mechanical properties (DMA) Dynamic mechanical analysis (DMA) provides a sensitive method to study the effects of thermal and morphological transitions on the modulus and the dynamic loss behavior of multi-phase polymers. A typical DMA graph of an SBCP with a low Tg and Tm of the SS and monodisperse crystallizable HS segments is given in Fig. 36. As can be observed, there was a low Tg of the SS, a high flow transition temperature Tflow due to the melting of the HS, a high modulus at very low temperatures (−100 ◦ C), a lower modulus above the Tg of the SS phase, which was only little temperature-dependent, and a low loss modulus. At −100 ◦ C, the shear storage modulus for polymers is generally on the order of 1100 MPa (E-modulus ∼ 3300 MPa). This low temperature modulus depends very little on the type and content of the HS. The SS Tg (maximum of the loss modulus) is dependent on the type and length of the SS (and any crosslink density) as well as on the amount of dissolved HS [70]. The modulus above the Tg of the SS, in the rubbery plateau region, depends on the reinforcing effect of the HS on the soft matrix. As the MD-HS crystallites have a ribbon structure with a high aspect ratio, the reinforcing effect is strong [26,70,71,77,78,81,83,126]. The tempera-

Fig. 36. Storage modulus G (solid line) and loss modulus G (dashed line) as functions of temperature from DMA measurements of PTMO2000 –T6T6T.

3.8.1. Tg soft segment The Tg of the soft phase depends on the chain flexibility, which in turn is a function of the chemical structure, the length between network points and the amount of dissolved HS. Monodisperse soft segments exhibit a slightly lower Tg than their polydisperse counterparts [2,3,9,12,70,85]. The Tg of polyethers with similar SS lengths was found to decrease in the order PEO > PPO > PTMO (Fig. 18). This is due to the ether group concentration in the polyether, giving rise to a stronger dipole interaction [127]. Increasing the length of the SS leads to a decrease in the crosslink density and, consequently, decreasing Tg . This decrease followed the behavior of chemically crosslinked polymers [70]. With short MD-HS, the HS concentration was low (4–35%) and the crystallinity high (70–90%), and only a small amount was dissolved in the SS phase (1–10 wt%). Thus, SBCPs with MD-HS often present a lower Tg of the SS phase as opposed copolymers with a polydisperse HS distribution, such as the commercial segmented block copolymers. 3.8.2. Soft Segment crystalline phase If a readily crystallizable segment has been employed for the SS, e.g., PTMO, PEO, PHA or PCl, then the SS melting temperature and crystallinity increase with an increasing SS length. The SS crystallites increase the modulus in the region between their Tg and Tm (Tflex ) and, as a result, the rubber plateau starts at Tm of the SS (Figs. 2 and 36). The SS melting temperature is often above room temperature and, if present, it affects the room-temperature properties. 3.8.3. Modulus The modulus in the region of the rubber plateau is a function of the SS crosslink density and reinforcement by the dispersed HS phase. With crystallizable HS, the modulus increased significantly with the crystalline HS content [26,68,70,71,74,75,78,81,84,126]. Parameters in this reinforcing mechanism include the stiffness of the fibers, the contents and the aspect ratio. Moreover, a useful model for this is the Halpin–Tsai relationship [70]. As with MD-HS, the crystallinity was elevated and the crystallites displayed a ribbon-like structure with a high aspect ratio (1000). Consequently, a strong increase in modulus with crystalline HS concentration was observed (Fig. 37). The G storage moduli increased with the HS crystalline content, and a similar increase was observed for the T6T, T6T6T and T6T6T6T segments. The modulus is thus not dependent on the HS length (i.e., the crystallite thickness), but sensitive to the HS crystalline content. For copolymers with less regular urethanes (e.g., TDI, MDI), the modulus increase with HS concentration is often lower [55].

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 Fig. 37. The storage modulus (G25 ◦ C ) of –(PPO–T(6T)x )n – as a function of the crystallized amide content. x = 䊉, 1 (diamide); , 2 (tetra-amide); , 3 (hexa-amide) (dashed line: Halpin–Tsai, L/D = 1000). Reproduced with permission from ref. [70], copyright (2006) Wiley & Sons.

Increasing the HS polydispersity was found to decrease the plateau modulus (Fig. 3). This was likely due to the lower HS crystallinity, and possibly also to a lower aspect ratio of the crystalline nano-ribbons [85]. The type of SS (PTMO, PPO, PEO, PEB, PDMS) had an effect on the modulus, and the moduli were particularly low for PEB and PDMS matrix segments [28,43,71]. It may be noted that an estimate of the volume of the crystalline content in a copolymer may be made from the plateau modulus. 3.8.4. Temperature dependence of the rubber plateau modulus The modulus of semi-crystalline polymers in the range intermediate to Tg and Tm decreases with increasing temperature due to the gradual melting of the crystallites (Fig. 22) [8,9]. Such a decrease in modulus with temperature is quantified as G (or E ) (%/◦ C) (Fig. 37), and for semi-crystalline polymers such as polypropylene, polyamides and polyesters, this decrease in modulus with temperature has been seen to be quite strong (2.3–16%/◦ C). Such a strong decrease was also the case for SBCPs with a polydisperse length HS distribution (1.3–4.3%/◦ C) (Fig. 3). However, the changes in modulus with temperature were low for the MD-HS materials, suggesting that all the crystallites melted at a single temperature. With certain MD-HS copolymers, even an increase in modulus with increasing temperature has been observed [26,70,71,79,84]. This resembled the behavior of an ideal rubber with a constant crosslink density with temperature, for which the modulus increase is due to an entropy elasticity effect [127]. With truly monodisperse HS, the crystallite thickness is constant and no partial crystallized interphase is thought to exist between the HS and the SS (Figs. 1 and 23). Deuteron NMR has been used to study the mobility in the MD-HS and no gradient in the MD-HS mobility was observed from the interface to the center of the HS [20]. Neither was there evidence of HS dissolution with temperature at temperatures lower than the melting temperature. However, it was possible that some irregular packing was present in the width direction (b) of the crystalline

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ribbons. A gradual melting of the partial crystallized interphase lowered the crystallinity with temperature, and consequently also the modulus. Such gradual melting can be studied with temperature-modulated FTIR [26,71,80] and SAXS [26,83] (Figs. 27 and 35). However, whether the HS was H-bonded or not did not seem to have any effect on this modulus–temperature behavior [8,9,83]. With an increasing HS polydispersity from 1.0 to 1.2, G was found to change from 0.3 to 1.1%/◦ C suggesting that for the 1.2 PDI sample, the crystallinity decreased more significantly with temperature [83]. One can say that SBCPs presented a slightly temperature-dependent modulus when the crystallites were uniform in thickness. This also suggests the HS crystallites are not uniform in thickness, and the HS are probably not monodisperse in length, for copolymers that have a temperature-dependent modulus in the rubber plateau region. 3.8.5. Hard segment Tg A hard segment Tg is present if an HS amorphous phase is present. If the HS and the SS form a single phase in the melt, generally no liquid–liquid demixing occurs upon cooling, and no HS amorphous phase is present. With Hbonded HS, the compatibility with the SS is often low, and only very short segments are able to mix. With the polyether–amide system, phase separation has been found to take place for an HS length of only three repeat units (six amide units) [14,70]. If the HS had a polydisperse distribution, liquid–liquid demixing would occur with an average value of only two repeat units [14,70]. This is similar to the behavior of the polyurethane and polyureas. With short MD-HD, such liquid–liquid demixing can be avoided, while at the same time the crystalline content can be high. 3.8.6. Flow temperature For segmented block copolymers with crystallized HS the melt transition (Tflow ) can be observed with DMA even at low HS concentrations (Fig. 36). The flow temperature (Tflow ) measured by DMA usually corresponds well with Tm measured by DSC [26,68,70–73,77,78,81,84]. The effect of using MD-HS on Tflow is the same as that on Tm by DSC, as discussed above. 3.8.7. G loss modulus The change in storage modulus with temperature can clearly be observed in DMA thermograms. For an accurate measurement of the transition temperatures, one generally employs loss modulus data (G ), and more specifically the temperature at the maximum of the loss modulus peak (Fig. 37). Also the amount of damping can be determined from the loss modulus–temperature curves, and a low loss modulus indicates a low damping loss at that temperature. The SBCPs with MD-HS have very low loss moduli in the rubber region (i.e., 0.3–1 MPa) and these values have been found to be little dependent on the storage modulus [26,68,70–72,77,78,81,84]. With increasing PDI (1.0–1.2), G increases (0.25–1.0 MPa) [85]. The SBCPs with MD-HS thus display a very low damping loss and exhibit excellent dynamic properties. Frequently, in DMA analysis, often tan ı = G /G is reported instead of G itself.

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Fig. 38. Schematic diagrams of the structure in the melt of (a) a multi-block and tri-block (–(AB)n –) copolymer and (b) a tri-block (BAB) copolymer. Reproduced with permission from ref. [110], copyright (2009) Wiley & Sons.

3.9. Melt rheology The melt rheological behavior of high-molecularweight SBCPs was similar to that of homopolymers, e.g., at low molecular weights (<10,000 g/mol) they presented a Newtonian behavior, and at high molecular weights they demonstrate a fairly strong shear rate dependence [27,110]. However, low molecular weight di- and tri-block copolymers with MD-HS at the ends had very high melt viscosities at low shear rates, and a strong shear thinning effect [30,110,128,129]. At high shear rates, on the other hand, the di- and tri-block copolymers presented melt viscosities similar to those of multi-block copolymers. After stopping the shearing melt thickened rapidly, demonstrating a thixotropic behavior. The di- and tri-block copolymers were thought to form a gel in the melt. The structure could easily be ruptured during straining, but reformed again after the strain was released. The HS end-blocks seemed to cluster in the melt (Fig. 38b), while the multi-block copolymer (Fig. 38a) did to have this behavior. 3.10. Tensile behavior A tensile test is easily performed, but for a system with two or more phases, the results are difficult to understand. The phase-separated hard phase segments have two functions: they act as physical crosslinks and at the same time as reinforcing fillers. The SBCPs have a yield point and, upon straining of a sample, the HS crystallites are sheared and oriented. The SS also orient and may strain-crystallize [37,86,130,131]. Studying the tensile properties on well defined SBCPs with MD-HS gives new insights into the deformation process.

tion [30,77,130,131]. Necking has been observed for the higher HS concentrations [30], and this yielding has been explained as a break-up of a percolation morphology (network structure) [30,48], and also as a shearing of long crystalline ribbons [80,130,131]. If no HS are present then the linear polymers are a viscous liquid with no yield point. The yield stress for copolymers with MD-HS of constant length was found to increase linearly with the HS content (Fig. 39) [77,130,131]. However, this increase with the HS concentration was stronger for longer HS (T6T6T) than for shorter HS (T6T) (Fig. 40). The thin, the non-H-bonded and polydisperse crystalline ribbons [85] exhibited lower yield stresses and were thus less resistant to shearing of the crystallites. The yield stress decreased with temperature according to the Eyring relationship [131]. Surprisingly, the modulus as a function of the amide concentration was not dependent on the HS length (Fig. 37), contrary to the behavior with the yield strength (Fig. 40). Also, the yield stress values displayed a temperature dependence while the modulus did not [131]. At the yield point, deformation of the crystallites had already begun, and the deformation of thin crystallites occurred more readily at higher temperatures. For semi-crystalline polymers, the modulus–yield strength relationship has been a matter of discussion for some time, and this topic has

3.10.1. Young’s modulus The Young modulus was found to increase with the MDHS content and this increase was similar to that of the shear modulus (G ) (Fig. 37) [70,77,78,130,131]. The E/G ratio was often a factor of 3, as expected for an effectively “incompressible material”, for which the bulk modulus is much greater than the shear modulus, often seen for polymeric materials in the rubbery plateau region. 3.10.2. Yield behavior These copolymers present a yield behavior, and the yield point is regarded as the onset of massive plastic deforma-

Fig. 39. Yield properties of PTMOx –T6A6T copolymers as function of HS concentration: , yield stress; , yield strain. Reproduced with permission from ref. [131], copyright (2008) Springer.

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Upon straining of SBCPs, particularly above the yield point, the HS crystallites become deformed. In their length (c), the nano-ribbons are held together by H-bonding which is not as strong as covalent bonds. Upon straining, these nano-ribbons are sheared and thereby broken up, and as a result, their aspect ratio is lowered [7,26,30,130,131]. Consequently, upon deformation, the amorphous phase is oriented, the crystalline phase is sheared as well as oriented, and strain induced crystallization of the SS may take place.

Fig. 40. Yield stress of –(PPO–T(6T)x )n – as a function of the amide concentration for various segment lengths x: 䊉, 1 (diamide); , 2 (tetra-amide). Reproduced with permission from ref. [70], copyright (2006) Elsevier.

been particularly studied for polyethylene [132]. However, for polyethylene, it is not possible to alter the crystallinity without changing the crystallite dimensions. The results given above showed that there is no unique relationship between modulus and yield strength. The yield strain was determined by the Considiére method and was found to be in the order of 25–70% [70,77,130,131]. The yield strain decreased with increasing MD-HS content (Fig. 39), reflecting a more rapid increase in the modulus than in the yield stress with increasing MDHS content. The yield strain also decreased when raising the temperature [131]. With increasing temperature, the modulus remained constant; however, the yield strength and consequently the yield strain decreased. The variations in yield strain values were thus related to changes in the ease of shearing the HS. 3.10.3. Past the yield point The deformation at higher strains has been found to depend strongly on whether strain hardening can take place. In SBCPs with MD-HS, the HS crystallinity is already high and strain crystallization of the HS can barely take place. If no strain hardening in the matrix phase occurs, the engineering stress remains constant with the strain while the fracture strain can still be high. However, if strain crystallization of the matrix phase is possible, as for PTMO SS, a significant strain-hardening effect has been observed at strains above 200% [26,30,77,80,130,131]. The strainhardening effect has also been shown to strongly increase with the molecular weight of the copolymer [77,130,131]. However, no strain hardening was observed when an SBCPs with PTMO SS was tested above the PTMO melting temperature [130,131]. This strain hardening in the copolymers was clearly due to the strain induced crystallization of the SS. A very unusual effect is the increase in SS melting temperature on orientation, e.g., a copolymer with PTMO2000 had a melting temperature of approximately 0 ◦ C when unoriented and of 43 ◦ C after orientation [130]. Thus, at room temperature, unoriented PTMO2000 copolymers at low strains exhibited a decent elastic behavior combined with high fracture properties due to strain hardening. Probably other crystallizing SS will also show this behavior.

3.10.4. Structural changes upon drawing It has been observed that, upon straining to 200%, a significant softening takes place and the modulus is lowered by a factor 7 (Fig. 41) [130]. Straining up to 200% led to a decrease in modulus despite the expected orientation of the nano-ribbon HS. As the modulus of an SBCP is mainly dependent on the reinforcement of the nano-ribbons, the sheared nano-ribbons were expected to have lower aspect ratios. Such a shortening of the nano-ribbon aspect ratio has been observed with AFM [30,37,122]. At high strains (>200%), a certain increase in modulus was seen (Fig. 41). This increase could partly be due to the orientation of the sheared nano crystals and partly also to the strain-hardening effect of the SS phase. The onset of strain softening has even been found to take place even before the yield point was reached [130]. Information on the orientation and structural changes of the HS crystallites in segmented block copolymers during straining can be obtained with WAXS [7,26,50,51], SAXS [7,26,37,50,51,119], birefringence [6,37,133], AFM [30,37,120,122] and IR-dichroism [6,26,86,133]. In the interpretation of the observations it is often assumed that the HS crystallites have a lamellar (plate like) morphology, with well stacked lamellae. However, such interpretations may be challenged since a poorly ordered nano-ribbon morphology is often observed with short monodisperse HS (Fig. 11). The IR-dichroism technique provides data on the orientation at the atomic scale and is informative for the deformation studies. IR-dichroism studies have been performed on polyurethanes [6,133], polyureas [26] and

Fig. 41. Strain softening: change of E-modulus upon straining of PTMO1000 –TT. Reproduced with permission from ref. [130], copyright (2001) Elsevier.

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Fig. 42. Orientation functions of various groups of the PTMO1000 –BDI diurea copolymer as studied by infrared linear dichroism as a function of strain and subsequent relaxation. Reproduced with permission from ref. [26], copyright (2006) American Chemical Society.

polyether–amides [86]. SBCPs with urea and amide MDHS are very suitable for this as they have strong crystalline HS peaks in the IR spectrum. Upon straining SBCPs with MD-HS, very peculiar things happen, as illustrated in the deformation of a PTMO–diurea copolymer (Fig. 42). The SS ether groups were only minimally oriented, even to strains as high as 700%, as the orientation function was only 0.15. On relieving the strain, the SS orientation was lost. If, however, the SS were to strain-crystallize, the ether orientation was stronger and the orientation could no longer be fully relaxed upon relieving the strain [86]. With PTMO2900 , the strain-hardening effect was evident with strains higher than 200%. Upon straining, the crystalline PTMO ribbons were thus more significantly oriented than the PTMO segments in the amorphous phase. The orientation of the HS was studied using the crystalline carbonyl and N–H bands [26,86,133]. At first, HS segments in the chains (c-direction) were oriented perpendicular to the strain direction and with a maximum orientation value at a strain of approximately 100% (Fig. 42). At higher strains, the orientation of the crystallites changed by approximately 90◦ as the HS chains (c-direction) were pulled in the drawing direction. Upon relieving the strain, the orientation only changed minimally. It has been observed with AFM that not only did the HS orientation change upon drawing but the crystalline ribbons were shortened and could no longer be observed on a highly strained sample [39,120]. 2D-SAXS analysis was carried out after pre-straining and relaxing. In the prestrain region of 150–500%, with SAXS a four-point pattern was observed, and at higher strain, a two-point pattern was seen. The four-point pattern was attributed to a tilted ordering of the HS segments [26,37]. The two-spot pattern at very high strains indicated the presence of rather small, fragmented crystalline domains that were ordered in stacks with correlations among them only in the direction parallel to the deformation axis. The dimensions of the long a-axis

of the crystalline ribbons were now expected to be only ∼5 nm [7] – a value much smaller than the >300 nm of the unoriented sample. All these studies show that the morphological changes of the HS crystallites upon straining are in steps, as depicted in Fig. 43. These studies showed three stages in the HS orientation. In the unoriented state, long dispersed crystalline ribbons are present with a random orientation (Fig. 43a). At low strains (50–200%), the long ribbons have been sheared to a certain extent and oriented with the ribbon length (a-axis) in the drawing direction (Fig. 43b). The orientation is not ideal and in the SAXS a four point pattern is seen due to a tilt of the nano-ribbons. At this stage the modulus was already considerably decreased (Fig. 41) [130]. Upon further straining (ε > 700%), the HS are deformed even more and their orientation is changed by 90◦ (Fig. 43c) compared to the 50–200% strain situation (Fig. 43b). The HS chains (c-axis) then lie in the drawing direction and a two-point SAXS pattern was observed. This rotation of orientation can only be possible if the length of the HS ribbons (a-axis) decreases to the size of the HS thickness (c-axis). The aspect ratio (a/c) was thereby expected to be lowered from >300 to ∼1. At the same time the HS nanoribbon length decreased [7,37,122], the HS strain and the SS strain-crystallized [7,37,133]. 3.10.5. Fracture behavior An SBCP can have high fracture strains and stresses. With increasing HS concentration, the fracture stress  frac increases, and the fracture strain εfract is reduced. The values of the fracture stress and strain can be combined into a single parameter: the true fracture stress ( true =  fract (1 + εfract )). This true fracture stress has been found to depend only weakly on the HS concentration or the HS polydispersity, but strongly on the molecular weight of the copolymer (Fig. 44) [72,77,85,130]. The true fracture stress has also been found to significantly increase with the strain-hardening behavior of the SS 77. The strain-hardening stabilizes the deformation and prevents an early fracture due to localized deformations. The high true fracture stresses approach the values for spider silk and industrial fibers [134,135]. SBCPs with SS that do not strain-crystallize have not only a lower fracture stress, but also a lower fracture strain. Also SBCPs with crystallizable SS have this behavior above the melting temperature of the SS chain [86,131]. Increasing the test temperature did not increase the fracture strain of an PTMO-based SBC as might be expected. Rather, it was decreased. With the absence of strain hardening at higher temperatures, the stabilization of the deformation is lost and the true fracture stress decreases with temperature. BAB tri-block copolymers with “sticky” groups at the ends had true fracture stresses as a function of the molecular weight that were similar to those of multi-block copolymers [72]. The fracture stresses and strains seemed to be particularly high with very short diamide and diurea segments [12,26,30,56,130]. These high fracture strains (>1500%) were surprising as the maximum draw ratio of a crosslinked network is about 600%. This suggests that,

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Fig. 43. Orientation and deformation of the nano-ribbon crystallites of segmented block copolymers upon straining. Reproduced with permission from ref. [130], copyright (2001) Elsevier.

during deformation of the SBCP, the short HS crystallites (network points) were easily sheared and thus allowed a higher maximum draw ratio than possible with a crosslinked rubber. The fracture toughness values of SBCPs are sometimes compared with other materials. Such a comparison is only meaningful if either the engineering fracture stresses or fracture strains are the same. If the true fracture stress is a useful parameter {fracture strain + 1) × fracture stress)}, then it is not possible that at the same time the fracture energy – which is approximately fracture strain × fracture stress/2 – is meaningful.

The CS values were found to increase with increasing PDI (Fig. 46) [85]. Thus, in the presence of some very thin crystallites, the elastic properties were lowered; probably because these thin crystallites were more easily deformed. The CS values for non-H-bonded HS samples were also higher and this too was believed to be due to their more easily deformable crystallites.

3.11. Elastic properties 3.11.1. Compression set A standard method for investigating elastic properties is the compression set technique. Often used conditions include 25% compression, for 24 or 48 h, at room temperature or higher (e.g., 70 ◦ C) and a relaxation time of 30 min after releasing the strain. A strain of 25% is lower than the yield strain for SBCPs as measured in tensile. The compression set (CS) values generally increase with HS content (increasing modulus) [30,79,129,130,136] and were found to be higher for the non-H-bonding HS [83]. It was surprising that the CS decreased with increasing length x of the MD segment, even though the modulus increased (Fig. 45) [70,78].

Fig. 45. Compression set as a function of the modulus for –(PPO–T(6T)x )n – with different HS lengths: 䊉, diamide (m = 1); , tetra-amide (m = 2); , hexa-amide (m = 3). Reproduced with permission from ref. [70], copyright (2006) Wiley & Sons.

Fig. 44. The true fracture stress of PTMOx –T6T6T copolymers as a function of inherent viscosities with varying T6T6T concentrations. Reproduced with permission from ref. [77], copyright (2003) Elsevier.

Fig. 46. Compression set results as a function of the polydispersity of PTMO2000 –T6A6T. Reproduced with permission from ref. [85], copyright (2006) Wiley & Sons.

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3.11.2. Tensile set Measurement of the tensile set TS can be used on a tensile bar or a fiber to study the elastic properties at higher strains. In a TS test, the sample is loaded, unloaded and reloaded without a standing time in between. The difference in the onset of strain between the first and second loading is taken as the set. The TS values for SBCPs with MD-HS segments have been found to be relatively low and increased with an increasing modulus (HS concentration) [26,70,78,136–138]. Several studies have been devoted to the TS behavior with cyclic loadings in which the strain is increased in steps and the set is determined after each strain step [26,70,78,136,138]. The TS as a function of the strain step gives interesting insights into the effect of strain on the TS (Fig. 47) [70]. Several regions can be observed. Up to the yield strain (∼40%), the TS seemed to increase linearly with the strain. At higher strains, a plateau region was observed and above 100% stain, the TS of PPO-T6T6T increased. Like observed with the CS (Fig. 45) the TS values decreased with increasing HS length x. Also with an increasing SS length, H-bonding and a decreasing PDI the TS values were lowered in the 0–100% strain region, as in the CS test. With an easily strain-hardening SS, the elastic properties up to 100% strain can still be excellent [26,136,138]. In a TS test, also the hysteresis of the deformation energy can be measured as a function of strain. Such graphs are relatively simple. For PPO-T6T6T, the hysteresis energy was independent of the applied strain when the latter was above 10% [70]. For PTMO–diurea, the hysteresis energy only stabilized above 200–300% [26]. 3.11.3. Stress relaxation SBCPs are viscoelastic materials and their timedependent properties are often studied in creep or stress relaxation (SR) experiments. Stress relaxation processes are Boltzmann processes, which are usefully expressed as functions of the logarithm of time owing to the wide range time-span of the response [139]. Elastic SBCP stress relaxation experiments can be carried out over a wide strain

Fig. 48. Stress relaxation of PTMO2000 –T6A6T at various strains [139]. Reproduced with permission from ref. [136], copyright (2009) Springer.

range, from below the yield strain to very high values. Studies have been performed on SBCPs with MD-HS at different strains, and the resultant stresses have been plotted on a log time scale (Boltzmann processes) (Fig. 48) [59,107,136]. Already during the first seconds, strong changes were found to take place and after approximately 100 s, the SR process became stabilized and the stresses decreased linearly with log time. The stress relaxation in the first 100 s increased with the strain but this value was influenced by the time of loading. For the time interval from 100 to 10,000 s, the slope of the curve was determined and, for all strains, this slope presented a 6% change for each time decade. The time interval for this SR might even start much earlier than 100 s, but the loading time interferes with accurate time measurements at very short times and possibly the two processes are just one. The SR was also studied when increasing the MD-HS concentration at a 25% strain [59,136]. The initial decay decreased with decreasing HS and this was probably due to the lowering of the modulus. However, the SR for the time interval from 100 to 10,000 s was practically independent of the HS concentration and had a value of about 6% per decade of time. Similar results were obtained for polyurea samples at 10% strain [30]. For PTMO1100 , the SR values were about 6% per decade of time whereas the corresponding value for PTMO2450 was 8%. These are low stress relaxation values. 3.12. Fiber properties

Fig. 47. The incremental tensile set as a function of strain as determined in a multi cycle test was studied on PPO2300 –T(6T)x – with varying HS length x: 䊉, diamide (x = 1); , tetra-amide (x = 2); , hexa-amide (x = 3). Reproduced with permission from ref. [70], copyright (2006) Wiley & Sons.

SBCPs with very short MD-HS present high fracture strains (>1500%) and the tensile stresses become elevated if strain hardening of the SS can take place. The elastic properties improve with increasing MD-HS length (Figs. 45 and 47). These materials can be used in elastic fiber applications. It would be extremely interesting if these materials could be melt-spun at high spinning speeds. For this, the HS crystallization rate must be high, which can be obtained by employing MD-HS. Upon drawing of these copolymers, the fracture stresses were found to increase while at the same time the fracture strains decreased. However, the true fracture stress remained con-

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stant, while the elastic properties improved somewhat on drawing [137,138]. Moreover, on melt-spinning SBCPs with MD-HS, strong spinline drawing was observed with true fracture stresses similar to those of post-drawn samples [137,138]. This suggests that these fast-crystallizing SBCPs can probably be processed with the simple meltspin process and thereby produce highly oriented elastic fibers with decent properties. The true fracture stresses observed for these soft SBCPs were found to be incredibly high (900 MPa) [130]. These elastic fibers thus combine good tensile properties with a high elasticity. Fibers that have a high elasticity at a higher modulus are very interesting as with a higher modulus they stick less together on the bobbin, which can be a major problem with soft materials. Copolymers with longer HS which combine a good elasticity with a higher modulus thus present an advantage, as shown in Fig. 47. The fiber forming PTMO–T6T6T copolymers have a high melting temperature, are melt processable, crystallize very fast on cooling and have low TS values and a relatively high modulus. Copolymers for which the matrix phase has a high Tm , such as a crystallizable polyester, can exhibit two high Tm s if modified with a high melting HS (Fig. 2d). Thus, drawing such a modified polyester is possible at a higher temperature, to a higher draw ratio and to give rise to a higher fracture strength of the fiber [105]. Also the crystallization temperature of the polyester increases upon drawing due to the presence of the HS crystallites (T6T6T). 3.13. Solvent resistance The solvent resistance of semi-crystalline polymers is superior to that of their amorphous counterparts. Segmented block copolymers with a crystalline HS phase are difficult to dissolve and thus have an improved solvent resistance. Particularly critical is the solvent resistance of amorphous polymers of high Tg . Studies have been performed to determine whether the solvent resistance of high-Tg amorphous polymers could be improved by copolymerization with MD-HS with a high melting temperature. As we have seen above, the melting temperature

Fig. 49. Solvent resistance of polysulphone–T6T6T copolymers of injection-molded samples suspended one hour in chloroform. Reproduced with permission from ref. [101], copyright (2009) Wiley & Sons.

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of a HS crystalline phase in these copolymers is also dependent on the type and concentration of the matrix segment (Fig. 25). Thus, the crystallization window decreases with decreasing HS concentration. Poly(phenylene ether) [96–99], polycarbonate [100] and polysulphone [101] have been investigated as highTg matrix segments and tetra-amide (T6T6T) as MD-HS. It was found that in SBCPs with a high Tg matrix material, the T6T6T segments could still crystallize in an injection molding process. Moreover, the dimensional stability of these materials was improved (Fig. 2c). For the monodisperse T6T6T segments, a crystallization window of about 70 ◦ C seemed to be sufficient for fast crystallization. Remarkably, the semi-crystalline copolymers presented an enhanced solvent resistance already at a very low T6T6T concentration (5 wt%) (Fig. 49) [101]. 3.14. Hydrophilic copolymers Hydrophilic copolymers are interesting for membrane and biomedical applications. The matrix phase in these materials determines the surface and transport properties and the dispersed phase provides the material with its dimensional stability. With the use of short MD-HS that nearly fully crystallize, the matrix phase is rather pure and the molecular weights of the matrix phase segments can be higher. Consequently, block copolymers with long PEO segments display significant water absorption values (>100%) [84,93,140] and low contact angles with water (<30◦ ) [89]. Nevertheless, they still show good tensile properties in the wet state [141]. These copolymers also have a significant water vapor transport [90] and higher CO2 gas transport values than commercial polyether-amides [91]. The fact that these properties are improved is probably due to the very low amount of dissolved HS in the polyether phase. SBCPs based on PPO are less hydrophilic and demonstrate lower water absorption values (>20%) but surprisingly high CO2 transport values [140,142]. Thanks to the methylene side groups, the PPO has a high free volume and thus excellent transport properties. 4. Conclusions and outlook Short monodisperse HS in segmented copolymers presents a good miscibility with the matrix segment, a very fast crystallization and a very high crystallinity. The short monodisperse HS have ribbon-like crystallites with a uniform thickness, a very long length and a very high aspect ratio. These short monodisperse segments with H-bonding groups (urethane, urea, amide) and/or aromatic groups have a high melting temperature. The structure regularity, the length of the monodisperse HS and the number of H-bonding units have a strong effect on the melting transition. For a degradation-free processing of copolymers with a melting temperature higher than 170 ◦ C, amide HS are very suitable. Monodisperse HS have the advantage that at concentrations as low as a few wt% they are effective for increasing the dimensional and solvent stabilities of the matrix material. At the same time the matrix polymer properties are little affected by the presence of the nearly fully crystallized HS, which result in excellent low temperature

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