Block architecture influence on the structure and mechanical performance of drawn polyurethane elastomers

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Research ArticleReceived: 1 October 2013 Revised: 18 November 2013 Accepted article published: 5 December 2013 Published online in Wiley Online Library: 7 January 2014

(wileyonlinelibrary.com) DOI 10.1002/pi.4669

Block architecture influence on the structureand mechanical performance of drawnpolyurethane elastomersBorja Fernandez-d’Arlas,a Marian Corcuera,a James Runtb and ArantxaEceizaa∗

Abstract

Some natural biopolymers such as spider silk exhibit superb mechanical properties, characterised by their great toughness.Synthetic polyurethane (PU) copolymers also endow great toughness but lack silk’s stiffness and strength. The aim of this workwas to elucidate the role of segment block architectural features that influence PU stiffness and strength after cold drawing.For this purpose PUs with varied soft segment character, crystalline versus rubbery, as well as with different hard segmentchemistries, 4,4′-diphenylmethane diisocyanate/1,4-butanediol versus 1,6-hexamethylene diisocyanate/1,4-butanediol, weresynthesised by a two-step polymerisation method. We found that the architecture of both block segments has a dramaticinfluence on drawn PU mechanical performance, in which PUs with crystallisable soft segments and crystalline hard segmentsare shown to have a greater impact on developing stiffer and stronger materials.c© 2013 Society of Chemical Industry

Supporting information may be found in the online version of this article.

Keywords: polyurethane elastomers; block architecture; copolymers; supertough elastomers; polyurethanes; spider-silk mimicry

INTRODUCTIONAchievement of highly elastic materials, deformable but with high

yield and tensile strengths, has been a challenging goal.1–3 Thisclass of tough materials could find applications in different fieldssuch as regenerative biomedicine, civil engineering or robotics.The development of supertough materials has been hampereddue to the lack of knowledge underlying macromolecularstructure–property relationships. In this context spider silkhas become a model material combining extensibility and

strength.4–9 Researchers have concluded that spider silk’s superbcharacteristics under tension arise from the copolymeric nature ofits polypeptidic protein sequences consisting of crystalline alanine-rich segments and amorphous glycine-rich domains, in addition

to the oriented chain anisotropy caused by fibre drawing.6–9

Polyurethane (PU) elastomers share with silks their copolymericnature, consisting of hard segments linked to soft segments

in an alternating fashion.10–15 Thermodynamic incompatibilitybetween hard segments and soft segments as well as the stericand physicochemical properties of each of the segments15

lead to certain microphase separated structures, consisting ofhard segment domains (about 5–25 nm in diameter) dispersedwithin a soft phase consisting predominately of soft segments.16

Despite the multiblock nature of PU macromolecules, theirmechanical properties are very different from those of spidersilk, typically much more compliant, deformable and weaker,17,18

and are highly dependent on their composition.19 Differentauthors have studied the morphology changes occurring withindifferent elastomeric PUs undergoing deformation. Bonart et al.used X-ray scattering to study the differences in morphology of

4,4’-diphenylmethane diisocyanate (MDI) based PUs synthesisedwith either ethylenediamine or hydrazine as chain extenders.20

They observed differences in interdomain spacing and also severeanisotropy of scattering patterns upon stretching, attributed tomacromolecular orientation along the stress direction. Kobersteinand Stein extended these studies using small-angle X-rayscattering (SAXS) to investigate the structure of PUs with differenthard segments built up by either MDI-1,4-butane diol (MDI-BD)or toluene diisocyanate and ethylene glycol.21 These studies alsorevealed differences in interdomain spacing as a function of hardsegment nature. More recently Yeh et al. studied the morphologyof low hard segment content poly(urethane)urea and concludedthat their system consisted of hard segments forming domainswith lamellar morphology with chain axis perpendicular to thelamellar axis. They also noted the possibility of other hard segmentsystems with higher polymerisation degree which would formhard fibrillar domains.22 Walezko et al. using SAXS and wide-angle

∗ Correspondence to: Arantxa Eceiza, Grupo ‘Materiales + Tecnologıas’(GMT), Departamento de Ingenier ıa Quımica y del Medio Ambiente,Escuela Politecnica, Universidad del Paıs Vasco/Euskal Herriko Unibertsitatea(UPV/EHU), Pza. Europa 1, 20018 Donostia − San Sebastian, Spain. E-mail:[email protected]

a Grupo ‘Materiales+Tecnologıas’ (GMT), Departamento de Ingenier ıa Quımicay del Medio Ambiente, Escuela Politecnica, Universidad del Paıs Vasco/EuskalHerriko Unibertsitatea (UPV/EHU), Pza. Europa 1, 20018 Donostia − SanSebastian, Spain

b Department of Materials Science and Engineering, Pennsylvania StateUniversity, University Park, PA, 16802, USA

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Block architecture effect on polyurethane elastomers www.soci.org

Table 1. Composition of hard segments (HS) and soft segments (SS) in the synthesised PUs along with the relevant transition temperatures for theshort segments and PU molecular weights

PU HS SS Mw (g mol−1) Mn (g mol−1) Mw/Mn

MDI-BD-(PHMC-b-PCL) MDI-BD PCL-b-PHMC-b-PCL (Tg = −80 ◦C) 76 800 41 600 1.84

MDI-BD-(PHMC) MDI-BD PHMC (Tm ≈ 50 ◦C) 56 900 37 600 1.51

HDI-BD-(PHMC) HDI-BD PHMC (Tm ≈ 50 ◦C) 68 160 43 000 1.58

Mw, weight-average molecular weight; Mn, number-average molecular weight; Mw/Mn, polydispersity index.

X-ray scattering (WAXS) studied the evolution of the morphologywith strain of three different PUs containing 1,6-hexamethylenediisocyanate (HDI) with BD as hard segments.23

In order to assess the best macromolecular structure todevelop mechanical performance close to that of spider silk fibre,we previously studied the influence of block morphology onthe tensile mechanical properties of oriented PU elastomers.18

We found that materials with hard segments with a glasstransition temperature below room temperature (Tg < T room)developed better macromolecular alignment than PU with glassyhard segments. Here we extend this research by studying themorphology and mechanical performance of PUs with hardsegments with a glass transition above room temperature(Tg > T room) and a different degree of crystallinity and chainstiffness, by comparing PUs synthesised with either MDI orHDI, and BD as a chain extender. Oligomeric polycarbonateand polycarbonate-polyester were incorporated as crystallineand rubbery soft segments, respectively. The chain architecturerequirements to achieve a supertough material after drawingthe PU macromolecules are discussed taking into accountexperimental results obtained from a variety of characterizationtechniques such as AFM, Fourier transform infrared spectroscopy(FTIR), WAXS, SAXS and mechanical testing.

EXPERIMENTALSynthesis of polyurethanesPUs were synthesised in a round bottomed five-neck reactor bythe common two-shot polymerisation technique, as reportedin previous publications.15,16,24 Briefly, in the first stage thepolyols were reacted at 95 ◦C for 5 h in the case of HDI(Bayer AG, Leverkusen, Germany) based PUs and at 60 ◦Cfor 2 h in the case of MDI (Bayer AG, Leverkusen, Germany)based PUs until a prepolymer was obtained. In the secondstage this prepolymer was reacted with the chain extender, BD(Sigma-Aldrich, Mardid, Spain), and the mixture was vigorouslystirred for 5–10 min until viscosity increase was appreciated.Then the mixture was poured into a poly(tetrafluoroethylene)

mould to complete polymerisation in a press at 100 ◦C for10 h. Poly(hexamethylene carbonate-co-caprolactone)diol (PCL-b-PHMC-b-PCL, 2010 g mol−1, Ravecarb 111, Polimeri Europa,Ravena, Italy) and poly(hexamethylene carbonate)diol, PHMC 1990g mol−1, Ravecarb 106 ( Polimeri Europa, Ravena, Italy) were usedas soft segments. The diisocyanate/polyol/chain extender molarratio was 3/1/2 in all copolymers. The chemical compositionof the studied PUs, along with main thermal transitions of theneat soft segments, synthesised PU molecular weights and PUabbreviations are gathered in Table 1. The molecular structure ofthe PUs is depicted in Fig. 1.

Materials drawing and tensile testingSamples for tensile testing were die cut from hot pressed films,prepared by heating to 180 ◦C for 5 min under pressure, followedby slow cooling to 30 ◦C. Specimens of ca 2.5 × 0.1 mm2 cross-section were drawn at 100 mm min−1 to different ratios λ usinga tensile machine (MTS-Insight10). The drawn ratios were λ =2, 3, 4, 5, 6, 7 and 8, where λ = L/L0 and L and L0 are thelength of the sample at a particular deformation and the initiallength of the sample, respectively. Samples were maintained atthose deformations for 10 min and during that time stress wasmonitored at 10, 20 30 s and 1, 2, 3, 5, 7 and 10 min. Tensile tests ofboth pristine and drawn samples were carried out at a crossheadspeed of 100 mm min−1 with an initial crosshead spacing of ca 8mm and with a load cell of 250 N. Sample thicknesses and widthswere measured again after drawing and then clamped in such away that only the drawn material experienced the tensile stress. Atleast three specimens were tested for each condition. Yield pointswere determined visually from the strain–stress curves taking theyield strain as the point close to the intercept of the tangent inthe elastic regime with the tangent in the pseudo-plastic regime.In cases where there was a clear drop in engineering stress afteryielding, the yield point was taken as the point where dσ /dε =0.25 For calculating true stress curves, the engineering stress σ wasconverted to true stress by multiplying by 1 + ε, where ε is thestrain (L – L0)/L0 or ε = λ − 1. Therefore, to a first approximation, it

Figure 1. Molecular structures of the synthesised PUs: (a) MDI-BD-(PHMC-b-PCL); (b) MDI-BD-(PHMC); (c) HDI-BD-(PHMC).

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Figure 2. Morphology of the studied PUs: 1 µm × 1 µm AFM phase images of as-cast (a) MDI-BD-(PHMC-b-PCL), (b) MDI-BD-(PHMC) and (c) HDI-BD-(PHMC).Phase angle scales are 15◦ , 60◦ and 100◦ for (a), (b) and (c), respectively. (d) Infrared spectra of the studied PUs.

was assumed that no volume change occurred when the materialswere stretched.

Nanostructure morphology characterisationAFM images were recorded in the tapping mode in air at roomtemperature using a Nanoscope IV, Dimension 3100 from DigitalInstruments. AFM samples were cast in poly(tetrafluoroethylene)4 × 4 cm2 base moulds from 1.5 mL of 50 mg mL−1 PU solutions intetrahydrofuran/dimethylformamide, THF/DMF (1:1), and solventswere evaporated with a temperature/pressure cycle of 2 h at 40◦C/800 ± 50 mbar and 24 h at 60 ◦C/600 ± 40 mbar. Films weredie cut and mounted over a rigid support for analysis.

Size exclusion chromatographyThe molecular weights and molecular weight distributions wereanalysed by SEC using a Perkin Elmer chromatograph equippedwith a binary pump and a refractive index detector. THF was usedas the mobile phase and the separation was carried out withinfour columns packed with particle gels with different nominalpore sizes. Elution rate was 1 mL min−1 and experiments wereperformed at room temperature. The samples were dissolved toca 1 wt% in THF. The molecular weights and molecular weightdistributions were based on a calibration curve created withmonodisperse polystyrene standards.

Infrared spectroscopyThe synthesised PUs were analysed by attenuated total reflectanceFTIR using a golden gate accessory device mounted over a Nicolet-Nexus FTIR spectrometer. Spectra were obtained by averaging 20scans between 4000 and 400 cm−1 with a resolution of 2 cm−1.

Differential scanning calorimetryDSC between −60 and 250 ◦C was performed on a Mettler Toledocalorimeter at a heating rate of 20 ◦C min−1 under constant N2

flow. Samples were analysed after being drawn to different ratiosand then relaxed at those strains for 10 min. Specimens of ca 4 mgwere cut and weighed in aluminium pans.

X-ray scatteringEvolution of the structure with strain from the molecular level tothe supramolecular scale was monitored by WAXS, which providedinformation on PU segment arrangements with respect to eachother, and by SAXS which was used to explore supramolecularlevel arrangements. WAXS patterns were measured on a RigakuDMAX/Rapid (λ = 0.154 nm at 50 kV and 40 mA). For analysingmaterials at specified elongations, an extension rig was usedas described elsewhere.26,27 The SAXS data were collected in aMolecular Metrology instrument with a CuKα source (λ = 0.154nm at 45 kV and 65 mA), using a two-dimensional multiwiredetector with a sample-to-detector distance of 1.5 m. Theaverage interdomain spacing 〈d〉 was calculated from azimuthallyintegrated SAXS data using the Bragg equation d = 2π/qmax,where qmax is the scattering vector of the maximum in theLorentz-corrected intensity q2I(q). All WAXS and SAXS patternsare displayed with the deformation direction oriented vertically.Consequently, the meridian is oriented parallel to the deformationdirection, and the azimuthal angle ψ = 0 was taken to coincidewith the equator.

Molecular structuresMolecular models were drawn using ACD/ChemSketch (Freewareversion) from ACD/Labs.

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RESULTS AND DISCUSSIONMaterials design and morphologyThe influence of soft segment crystallinity on the mechanical per-formance of the drawn samples was analysed previously.18 In thepresent study this influence is analysed by comparing two PUs con-taining more typical hard segments such as those formed by MDI-BD. The PUs were MDI-BD-(PHMC-b-PCL) and MDI-BD-(PHMC). Bycomparing the behaviour of MDI-BD-(PHMC) with HDI-BD-(PHMC)the influence of hard segment crystallinity on the mechanical prop-erties after drawing is studied. In previous studies18 we comparedmaterials with hard segments with either Tg,HS > T room (glassy,amorphous) or Tg,HS < T room (rubbery, semicrystalline). In this workthe comparison is made between PUs with different hard segmentcrystallinities, all with Tg,HS > T room.15 Differences in molecularstructure can be seen in Fig. 1. MDI based PUs tend to developmore phase mixed systems than HDI based PUs due to differencesbetween diisocyanate structures, differences in phase separationkinetics and segment crystallinity.15,28,29 It is known that whileMDI-BD hard segments exhibit a staggered conformation requi-site for hydrogen bonding,28 the flexible linear aliphatic nature ofHDI-BD hard segment systems is likely to arrange the urethanegroups lying in the same plane.29 The cyclic and aromatic natureof MDI may restrict formation of low entropy associated structuresdue to rotational steric barriers as well as to partial double bondsarising in the main polymeric backbone as a consequence of res-onance between the aromatic rings and the urethane groups.15

These factors make the HDI-BD segments more prone to crystalli-sation and phase separation15 as can also be observed in the AFMimages presented in Fig. 2. The PUs under investigation displaydifferences in self-assembly. From Figs 2(a)−2(c) differences ofordering are observed: from small worm-like structures in Fig. 2(a)to long and straight nanofibrils in Fig. 2(c). Figure 2(b) displays anintermediate structure between Figs 2(a) and 2(c), perhaps arisingfrom the appearance of fractions of crystallisable soft segmentstogether with MDI-BD worm-like structures. Another reason for thedifferences between the morphologies of MDI-BD-(PHMC-b-PCL)(Fig. 2(a)) and MDI-BD-(PHMC) (Fig. 2(b)) could be differencesof solubility of the MDI-BD segments in their correspondingsoft phases.

Differences in segment organisation were also observed by FTIRanalysis. As can be seen in the spectra shown in Fig. 2(d), HDIbased PU presents bands typical of highly phase separated PUsin which N−H and C=O groups are tightly hydrogen bonded.13,15

This is observed in the N−H and C=O stretching regions whereshoulders at 3365 cm−1, 1710 cm−1 and 1660 cm−1 appearin HDI based PU, corresponding to hydrogen bonded N−H,hydrogen bonded carbonyls (C=O) from soft segment carbonatesand hydrogen bonded urethane (−NH−CO−O−), respectively.15

These shoulders are not so prominent in MDI based PU because ofa more phase mixed state of the MDI-BD hard and soft segments.15

Other PU characteristic bands are those around 1600 cm−1 for MDIbased PU, corresponding to aromatic ring stretching, and thosearound 1530 and 1240 cm−1 corresponding to N−H bendingand C−O−C backbone stretching, common to all PUs. Anothercommon band in all PUs is that appearing at 750 cm−1 attributedto the ‘rocking’ of more than four methylene units. The lowfrequency carbonyl bands below 1650 cm−1 are usually attributedto carbonyl vibrations within compact β-sheets.30 The PU with thelowest frequency carbonyl band was HDI-BD-(PHMC) with a bandat 1663 cm−1, attributed to carbonyl present in the HDI-BD packedhard segments.

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Morphology evolution studied by DSCFigure 3(a) displays the DSC thermograms of MDI-BD-(PHMC-b-PCL) as a function of strain. Although the soft phase Tg isindependent of strain, a small decrease in hard phase Tg (Tg,HS)with increasing draw ratio is observed. This modest decrease maybe associated with some phase mixing of weakly inter-associatedMDI-BD units with PHMC-b-PCL arising from the deformation. Asseen in Fig. 3(b), on stretching MDI-BD-(PHMC) a transition (Tm,SS)appears and grows in intensity with increasing draw ratio. Tm,SS

corresponds to the melting of soft segment crystals induced bystrain. In systems containing MDI-BD hard segments, an increaseof the small endotherm (Tm2,HS) is observed with increasing strain,which can be attributed to some strain induced crystallisation ofMDI-BD hard segments (see Figs S1−S3). However, the enthalpyof the higher temperature endotherm (Tm1,HS) seems to bereduced at higher draw ratios. This could mean that MDI-BDlamellar domains break down. In systems with PHMC as thesoft segment (Figs 3(b) and 3(c)) the soft segment glass transitiontemperature Tg,SS is observed to increase after drawing, suggestingreduced mobility of soft segment blocks due to the strain inducedmacromolecular organisation and packing. In contrast, as noted


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(a)

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Figure 4. Two-dimensional WAXS patterns of (a) MDI-BD-(PHMC-b-PCL), (b) MDI-BD-(PHMC) and (c) HDI-BD-(PHMC) non-drawn and drawn at λ = 4.5,and two-dimensional SAXS patterns (right) of PU at λ = 4.5. (d) Average interdomain spacings at different deformations for the PUs, as obtained fromSAXS. The deformation direction is vertical for both SAXS and WAXS data, and the azimuthal angle ψ = 0 coincides with the equator.

above, Tg,SS did not change after drawing MDI-BD-(PHMC-b-PCL).This is probably related to the rubbery state of PHMC-b-PCL at roomtemperature that enables molecular reorganisation and relaxationafter straining.

Slightly higher hard segment crystallisation was attained in thecase of MDI-BD-(PHMC) than MDI-BD-(PHMC-b-PCL) after drawing(see Figs S1 and S2). The soft segment melting endothermsachieved after pre-drawing systems with PHMC were lower forMDI-BD-(PHMC) (5.7 J g−1 for λ = 6 and 10.9 J g−1 for λ = 8)than for HDI-BD-(PHMC) (12.5 J g−1 for λ = 6 and 15.0 J g−1 forλ = 8) and the peak of the melting endotherm was higher forthe HDI-BD-(PHMC) system (49 ◦C) than for the MDI-BD-(PHMC)system (43 ◦C), suggesting that in the case of HDI-BD-(PHMC)thicker and more regular soft segment crystals were attained afterdrawing. This was also confirmed by X-ray scattering.

Morphology evolution studied by X-ray scatteringBoth SAXS and WAXS patterns acquired in situ at different strainsrevealed a different structure evolution with strain as a function ofblock architecture. WAXS patterns of undeformed PUs and of PUsstretched to λ = 4.5 are presented in Figs 4(a)−4(c). SAXS patternsof materials deformed at λ = 4.5 are presented to the right of Figs4(a)−4(c). WAXS provides insight into the disposition of atomsor a small number of atom segments with respect to each other,in contrast to SAXS which gives information about larger lengthdomain organisation. In all cases a dramatic change in scatteringpattern is observed after drawing. Anisotropic WAXS patternsin which maximum intensity was observed at azimuthal anglesperpendicular to the stretching direction of aligned or stretchedfibres have been reported previously for PUs,11,31,32 polyureas,27

poly(urethane)ureas,22 polyethylene,33–35 and spider and worm

silk fibres.36–38 Similar scattering behaviour was observed in thestudied PU samples.

From integrated SAXS patterns it was observed that in allcases the average crystalline interdomain spacing 〈d〉 decreasedwith strain, as shown in Fig. 4(d). This can be related to threedifferent phenomena occurring simultaneously: (i) strain inducedcrystallisation,11 increasing the number of electron-rich domains,(ii) hard domain reordering, and perhaps to a small extent (iii)Poisson’s effect caused by deformation, which pulls polymerchains closer to one another. Reduction of the mean crystallineinterdomain spacing obtained from SAXS data was reported by

Waletzko et al.11 for three different PUs which could also beattributed to strain induced crystallisation of the polyethyleneoxide units of their PU soft segments and therefore a decrease inthe distance between high electron density domains. In contrastYeh et al.22 observed an increase in domain distance by SAXSwhen stretching a segmented poly(urethane-urea) elastomerwith semicrystalline polytetramethylene oxide soft segments.They attributed this to lamellar hard segment breakdown butit could also mean that in highly phase separated systems suchas poly(urethane-urea)s the hard interdomain distances, ratherthan strain induced soft segment crystallisation or hard segmentphase mixing, prevail in the electron density variances (overall highelectron density domain distances) observed at different strains.Here a small extent of strain induced crystallisation was observedin all PUs, including MDI-BD-(PHMC) hard segments, as discussedpreviously, but hard domain rearrangements are not discarded aswill be discussed next.

WAXS intensities as a function of scattering angle 2θ aregathered in Fig. 5 (note that the line scans in Fig. 5 displayscattering out to higher 2θ than the two-dimensional patterns inFig. 4). Non-deformed samples containing MDI-BD hard segmentspresented less sharp diffraction peaks at 2θ = 20◦−22◦, while inthe case of the PU with HDI-BD hard segments this angular regiondisplayed three sharp peaks at 20◦, 22◦ and 24◦ (peaks α, β andθ , respectively) with the highest intensity at 22◦, reflecting thehigher crystallinity of HDI-BD hard segments compared with thoseof MDI-BD. Low intensity reflections of the HDI-BD-(PHMC) samplebetween 35◦ and 50◦, seen in Fig. 5(d), are presumably relatedto soft segment crystals that are formed due to the good phaseseparation of HDI-BD segments in comparison with MDI-BD. Ascan be seen in Fig. 5(a) these reflections appear again grown inrelative intensity as meridional reflections in both HDI-BD-(PHMC)and MDI-BD-(PHMC) PUs as a consequence of PHMC soft segmentstrain induced crystallisation. The reflections disappeared whenthe strained samples were thermally treated by immersing them indistilled water at 80 ◦C for 5 min, reflecting that at this temperaturesoft segments melt and lead to randomly oriented crystals andthat these reflections can be attributed to partially broken andoriented soft segment crystallite lamellae (see Fig. S4). This wasalso the case for the reflections at ψ ∼ 33◦ and 2θ = 14.5◦

(peak δ) observed in the PHMC polyol based PUs. In particularthese last reflections could be attributed to different carbonates


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= 33◦ and (c) ψ = 0◦ (equator). (d) Profiles of non-deformed materials asobtained from isotropic integration.

situated in different crystallographic planes of the extendedmolecular chains. A detailed study of the crystal structure and celldetermination of hexamethylene carbonate segments is outsidethe scope of this work.

Figure 6 displays SAXS intensity as a function of azimuthal angleψ for materials strained at λ = 6. The difference between samplepatterns should be noted. Meridional scattering (parallel to strain)was intensified for materials containing PHMC as the soft segment,as can be seen in the peaks at ψ = 90◦ and ψ = 270◦ (peak ω). In thecase of the material with PHMC-b-PCL soft segments the intensityonly appeared to increase along the equator (perpendicular tothe strain direction). The scattering produced along the meridianin samples with PHMC might be due to the crystalline natureof the PHMC segments which upon stretching their lamellaeundergo partial breakdown and orientation. PUs with MDI-BD

Figure 6. Azimuthal angle (ψ ) intensity profiles of the two-dimensionalSAXS patterns corresponding to MDI-BD-(PHMC-b-PCL), MDI-BD-(PHMC)and HDI-BD-(PHMC) at λ = 4.5.

hard segments develop scattering at equatorial angles observedas peaks at ψ = 0◦ and ψ = 180◦ (peak π ), while in the caseof the PU with HDI-BD hard segments these peaks are not soprominent. PU MDI-BD-(PHMC) evolved with an intermediatescattering pattern between MDI-BD-(PHMC-b-PCL) and HDI-BD-(PHMC), with four peaks (0◦, 90◦, 180◦ and 270◦), but contrary toHDI-BD-(PHMC) with the most intense maxima at 0◦ and 180◦. It istherefore concluded that scattering at equatorial angles is highlydependent on the nature of the hard segment. Since small-anglescattering patterns are sensitive to electron density variances ona nanoscale (1–100 nm) the differences should be attributed todifferences in scattering vectors caused by different arrangementsand morphology of the crystalline domains (electron-rich domains)at different strains. The differences between PUs should alsodepend on their initial degree of phase separation. In weakly phaseseparated PUs like those containing MDI, with hard segments witha lamellar nature, strain could also cause hard segment straininduced reordering, phase mixing and crystallisation (see FigsS1 and S2 for more detailed calorimetric studies of structureevolution with strain). Some of the lamellae might undergocomplex morphology evolution which might be induced bythe competition between the continuum mechanical torqueand molecular mechanical torque that would cause lamellarbreakdown and orientation at different degrees.22 At intermediatestrains these phenomena might cause scattering in a perpendiculardirection due to lamellar orientation along the strain with chainsperpendicular to the strain. Contrarily, in highly phase separatedmaterials such as HDI-BD-(PHMC) with a more fibrous hard domainnature, continuum and molecular mechanical torques would beencompassed and hard domain integrity might be preserved evenat high strain levels, developing only rotation and alignment alongwith marginal hard segment lamellae breakdown (see Fig S3 fordetailed calorimetric study of the structural evolution with thestrain of HDI-BD-(PHMC) polyurethane).

Mechanical propertiesMaterials were drawn to different elongations and stress relaxationwas monitored at these strains for each PU over a range of 10 min.


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Figure 7. Stress−relaxation plots of stress as a function of time for different draw ratios: (a) MDI-BD-(PHMC-b-PCL); (b) MDI-BD-(PHMC); (c) HDI-BD-(PHMC).

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Table 2. Summary of the mechanical properties of pristinePU and PUs drawn to λ = 7

Material E (MPa) εy (%) σ y (MPa) εmax (%) σ max (MPa)

MDI-BD-(PHMC-b-PCL) 14 ± 2 74 ± 9 4 ± 1 892 ± 80 10 ± 1

MDI-BD-(PHMC-b-PCL)drawn 4 ± 1 786 ± 37 32 ± 2 892 ± 30 37 ± 2

MDI-BD-(PHMC) 21 ± 2 29 ± 4 4 ± 1 1185 ± 127 19 ± 1

MDI-BD-(PHMC)drawn 26 ± 7 264 ± 94 68 ± 5 295 ± 126 71 ± 4

HDI-BD-(PHMC) 18 ± 3 47 ± 7 5 ± 1 1626 ± 200 33 ± 1

HDI-BD-(PHMC)drawn 113 ± 25 90 ± 5 95 ± 12 137 ± 52 102 ± 12

The results of the experiments are shown in Figs 7(a)−7(c). Itis observed that the higher the crystalline fraction of soft andhard phases in the PU is, the higher the stress σ is over thestudied range. Comparing Figs 7(b) and 7(c) we observe that thehard segment nature greatly influences the relaxation kinetics ofthe macromolecular PU chains. PUs with MDI-BD segments relaxmore rapidly than those with HDI-BD units. This can be explainedby the strong intermolecular hydrogen bonding between HDI-BD segments,15 which act as crosslinks, preventing interchainslippage and therefore retaining the tension at the given strainrather than flowing. This does not necessarily involve better MDI-BD domain alignment along the stress direction than HDI-BD sincehard segment bulkiness might hinder segmental mobility andmacromolecular alignment,18 but they are more easily protrudedalong the stress direction.

Drawing the materials leads to a significant change in thetensile properties of the PUs. In Fig. 8 representative stress–straincurves before and after drawing to different ratios are shown.The observed changes are related to the chain architecture ofthe PU blocks, in that the more crystalline materials are morecapable of forming load bearing ordered structures perpendicularto the strain direction. As an example, Table 2 shows the tensilemechanical properties of both pristine and drawn PUs at λ = 7.After drawing, the yield stress σ y increased 800%, 1700% and1900% for MDI-BD-(PHMC-b-PCL), MDI-BD-(PHMC) and HDI-BD-(PHMC), respectively. This behaviour is attributed to a change inmorphology in which the hard and soft crystals are isotropicallydistributed within the bulk to a more fibrillar structure in whichhard and crystalline domains are able to withstand loads along thestrain direction due to their favoured orientation in this direction.The low strain mechanical properties of the three PUs drawn todifferent draw ratios are shown in Fig. 9. It can be seen in Fig. 9(d)that the elastic modulus E decreased for the three PUs for drawratios below λ= 3. Above λ= 3 drawing MDI-BD-(PHMC-b-PCL) ledalso to a decrease in modulus while drawing MDI-BD-(PHMC) led tono significant variation in modulus, and drawing HDI-BD-(PHMC)led to a significant increase in modulus. This trend may be related tothe strain induced phase mixing of MDI based PUs which destroysenthalpically driven inter-associations formed during cooling ofthe PU films,18,27 while the increase in modulus of HDI-BD basedPU can be attributed to both soft segment crystallisation and harddomain orientation along the stress direction. The yield strainεy increased for all materials up to draw ratios of ca 3. After a


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Figure 9. Summary of the mechanical properties in the low strain regimeof the three different PUs drawn to different elongations.

draw ratio of 3 yield strain continued to increase for MDI-BD-(PHMC-b-PCL) while it remained constant for MDI-BD-(PHMC) anddecreased for HDI-BD-(PHMC). Therefore the increase in resiliencethat is observed in Fig. 9(a) is due to different combinations ofyield strain and stress. In the case of MDI-BD-(PHMC-b-PCL) thehigh resilience of 75 ± 3 MJ m−3 is due to a large yield strain of790% ± 37% while in the case of HDI-BD-(PHMC) it is due to a largeyield stress of 95 ± 12 MPa.

In Fig. 10 the evolution of the mechanical properties at breakwith draw ratio for the three PUs is presented. The tensile strengthincreased for all PUs after drawing but the increase was largerwhen crystallisable block segments were incorporated into themacromolecular PU design. HDI-BD-(PHMC) exhibited a strengthof 120 ± 5 MPa when drawn to λ = 8. This is a remarkable valuefor a thermoplastic polymer. Unfortunately this seemed to be thelimit for this PU, since further pre-drawing did not increase thestrength. The strain to break, εmax, decreased as the draw ratioincreased. This is a typical phenomenon of other drawn polymerssuch as ultrahigh molecular weight polyethylene.35 The toughnessdata are somewhat scattered but present subtle changes. Thetoughness of HDI-BD-(PHMC) seemed to be reduced after drawing,which could be related to the higher crystallinity attained and tothe faster damage propagation through the material. The MDI-BD-(PHMC) PU seemed to present a toughness maximum for

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Figure 10. Summary of mechanical properties at break of the threedifferent PUs drawn to different elongation ratios.

materials drawn to intermediate ratios, λ = 4–6, while that of theMDI-(PHMC-b-PCL) based PU appeared to increase steadily by asmall amount at the studied draw ratios. The fact that toughnessdid not vary dramatically with drawing ratio could mean that theoverall cohesive energy of the material remained unperturbedregardless of the supramolecular arrangement of the chain and itsblocks. This supposition might become more relevant at slowerstrain rates, when macromolecules might approach a pseudo-equilibrium state at each differential strain, rearranging ultimatelyto similar equilibrium conformations in all cases.

The increase in stiffness and yield strength coincides withthe increase in segment crystallinity. The notable differencesbetween MDI-BD-(PHMC) and HDI-BD-(PHMC) indicate the greatimportance of the hard segment nature. The differences are relatedto differences in molecular conformation and in-plane orientationbetween urethane groups in the cases of MDI and HDI based hardsegments, which prevent or allow PU crystallisation, respectively.Hard segment crystalline lamella orientation perpendicular tothe stress direction with the chain axes in the crystals alignedtowards the stretch direction (see Fig. S5) might act as loaddistributing network points and will have an influence on thesuperior properties of HDI-BD based PU after drawing. Parallel


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Figure 11. (a) True stress–strain representative curves of MDI-BD-(PHMC-b-PCL), MDI-BD-(PHMC) and HDI-BD-(PHMC) PUs drawn to a ratio of λ = 8,along with those of Argiope argentata spider major ampullate and capture spiral silk fibres.5 Also shown is the curve for HDI-BD-(PHMC) drawn to λ = 5.(b) Absorbed energy with strain of the materials from (a).

results have been found in our previous work in which wecompared isophorone diisocyanate (IPDI) with HDI based PUs.18

In Fig. 3 we can also observe that even when drawing PUs basedon MDI the hard segment melting enthalpies were equal to that ofthe drawn HDI-BD (see Figs S1−S3). This is attributed to the poorcapability of MDI-BD to crystallise even under strain.

In Fig. 11 true stress−deformation curves of materials pre-stretched to draw ratios of λ = 8, in which a constant samplevolume between the clamps was considered to account for thereduction in the cross-sectional area, are presented along withcurves of both major ampullate and capture spiral fibre silksfrom the Argiope argentata spider.5 In the case of the drawnHDI-BD-(PHMC) it is observed that while the stiffness of the PUwas over one order of magnitude higher (ca 100 MPa for thedrawn PU versus ca 3 MPa for capture spiral) the strength ofthe PU was about 50% of that of the capture spiral (295 MPafor the drawn PU versus 534 MPa for the capture spiral). Still,the engineering strength of this material was remarkably highconsidering it also presented an engineering deformation as highas about 170%. To show these significant properties, in Fig. 11(b)the absorbed energy at different strains of materials pre-stretchedto λ = 8 is represented together with Argiope argentata silks. Itis seen that while major ampullate silk absorbs a large amountof energy with small deformation, capture spiral silk absorbsmore energy but with higher deformation. Materials with PHMChave an intermediate behaviour. For comparison in Fig. 11(b) arepresentative curve corresponding to HDI-BD-(PHMC) drawn toλ = 5 is also shown. It is seen that some specimens of these typesof PU had high toughness, absorbing energy with intermediatedeformation between major ampullate and capture spiral silks. Inprevious studies18 some PUs were observed with a true stress ofca 800 MPa and an elastic modulus of ca 98 MPa while retainingan elongation to break of ca 480%. These combined propertiesaccounted for toughness of over 400 MJ m−3. This high elasticitycombined with high strength was attributed to the high straininduced crystallinity of the soft phase, favoured by hard segmentflexibility with Tg,HS < T room.

While HDI-BD-(PHMC) displayed a greater degree of hydrogenbonding in companion FTIR studies than the other PUs under

investigation, hydrogen bonding units of spider silk might bemuch closer to one another than in any of the PUs, since theamorphous regions are mainly composed of short poly(glycine)units.6 This would result in a higher degree of hydrogen bonding inspider silk across the fibre axis than in the PUs, whose amorphousregions consist of oligomeric polyols. A higher degree of hydrogenbonding results in a more efficient load distribution and usually inhigher stiffness and strength. Therefore highly hydrogen bondedPUs with a high degree of macromolecular orientation couldpossibly approach the mechanical performance of silk fibres.

CONCLUSIONSThis work has shed light on the role that block nature has onthe cold-drawing mechanisms of segmented PUs. The nature ofthe segments strongly influences the material’s strain inducedcrystallisation and macromolecular orientation and packing. Amaterial with rubbery soft segments and less crystalline MDI-BDhard segments leads to lower modulus specimens after drawing,due presumably to lamellar MDI-BD phase mixing with rubberysoft segments. Rubbery soft segments consisting of blocks ofPHMC copolymerised with PCL do not undergo strain inducedcrystallisation. PUs based on crystallisable soft segments lead todrawing induced strengthening in both cases studied but juststiffening in the case of the PU with highly crystalline hardsegments, which can be attributed to better hard segmentorientation18 and to preserved interchain association betweenmore fibrous HDI-BD based hard segments. Yield strength andresilience were in all cases improved after drawing. Althoughstrength was increased after drawing, material toughness seemedto be conditioned by its chemical nature but did not dramaticallydepend on drawing ratio. These results show the importance of PUblock structure when pursuing stiff, elastic and strong materials.This and previous work18 indicate that the mechanical propertiesof capture spider silk can be mimicked by properly engineeredPU elastomers and at the same time points to the need for newsynthetic mechanisms and strategies for common PU chemistryto approach the outstanding stiffness, strength and toughness ofmajor ampullate fibres.


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ACKNOWLEDGEMENTSB. Fernandez-d’Arlas acknowledges the University of theBasque Country (UPV/EHU) post-doctoral grant Ayuda a laEspecializacion de Doctores for its financial support and allauthors acknowledge the General Research Services of theUniversity of the Basque Country (SGIker), and especially theMacroconducta-Mesoestructura-Nanotecnologıa unit, for theirtechnical support. The authors also acknowledge funding fromthe Basque Government through the SAIOTEK 11-S-PE11UN132programme and in the frame of Grupos Consolidados (IT-776-13).B. Fernandez-d’Arlas also acknowledges a UPV/EHU post-doctoralmobility grant for accomplishing a 3 months’ stay at Penn StateUniversity. The authors are very grateful to Dr Raegan Johnson, DrNichole Wonderling and Dr Autchara Pangon for their help withthe X-ray instrumentation and techniques. JR would like to thankthe National Science Foundation, Polymers Program, for supportof the portion of this research conducted at Penn State, undergrant DMR-1206571.

This work is dedicated to the memory of Professor Dr Inaki B.Mondragon (1954–2012) who founded the Grupo ‘Materiales +Tecnologıas’ in 1988. He was an inspiration for all of us.

SUPPORTING INFORMATIONSupporting information may be found in the online version of thisarticle.

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Block architecture influence on the structure and mechanical performance of drawn polyurethane elastomers