Semicrystalline Non-Isocyanate Polyhydroxyurethanes as

Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

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Semicrystalline Non-Isocyanate Polyhydroxyurethanes as Thermoplastics and Thermoplastic Elastomers and Their Use in 3D Printing by Fused Filament Fabrication Vitalij Schimpf,†,§ Johannes B. Max,† Benjamin Stolz,† Barbara Heck,‡ and Rolf Mülhaupt*,†,§

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†

Freiburg Materials Research Center (FMF) and Institute for Macromolecular Chemistry, University of Freiburg, Stefan-Meier-Strasse 21 and 31, 79104 Freiburg, Germany § JONAS − Joint Research on Advanced Materials and Systems, Advanced Materials & Systems Research, BASF SE, Carl-Bosch-Strasse 38, 67056 Ludwigshafen, Germany ‡ Institute of Physics, University of Freiburg, Hermann-Herder-Strasse 3, 79104 Freiburg, Germany S Supporting Information *

ABSTRACT: Chemical fixation of the greenhouse gas carbon dioxide with diepoxides followed by melt-phase polyaddition of the resulting difunctional cyclic carbonates with 1,12diaminododecane (DDA) yields semicrystalline polyhydroxurethane (PHU) thermoplastics. Also, 100% biobased semicrystalline PHU thermoplastics are feasible. Opposite to conventional polyurethane syntheses, neither isocyanates nor phosgene are required as intermediates. Preferably, melt-phase polyaddition is performed in a twin-screw compounder in the absence of catalysts, which also catalyze side-reactions. Calorimetric measurements and small-angle X-ray scattering reveal the fundamental structure−property relationships governing PHU crystallization. The PHU melting temperatures vary between 40 and 115 °C, and PHU Young’s moduli range from 220 to 1430 MPa. Moreover, non-isocyanate PHU thermoplastic elastomers (TPHE) are readily tailored via melt-phase polyaddition of diamine-terminated flexible PHU prepolymers serving as soft segments combined with semicrystalline PHU as hard segments. As verified by means of thermal analysis (DSC), dynamic mechanical analysis (DMA), X-ray diffraction (SAXS), and microscopy (AFM), the careful balance between soft and semicrystalline hard-segment incorporation accounts for nanophaseseparation, which in the past has failed as a result of phase intermixing resulting from strong hydrogen bonding between soft and hard segments. For the first time, tailored PHU thermoplastics are employed in extrusion-based additive manufacturing by means of fused deposition modeling (FDM) or fused filament fabrication (FFF). Clearly, the presence of hydroxyl groups and their hydrogen bonding improves filament fusion and adhesion essential for achieving mechanical properties similar to PHU melt extrusion without encountering warpage.

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INTRODUCTION Since the pioneering advances by Otto Bayer in 1947, polyurethanes have evolved as versatile polymeric materials with an annual production of around 18 Mt in 2016 accounting for 5% of the global plastics market.1,2 Tailormade polyurethane materials meet highly diversified market needs ranging from flexible foams for bed mattresses and automotive seats to hard foams for thermal insulation, coatings for surface protection, textile fibers, adhesives, and sealants.3−5 Today, commercial polyurethanes are prepared by a stepgrowth polymerization of polyols with polyfunctional isocyanates.5,6 Taking into account their high reactivity, moisture sensitivity, and toxicity, handling of isocyanates requires special safety precautions and adjusted manufacturing procedures.7−10 At the beginning of the 21st century, the increasing social demands for sustainability, low carbon footprint, bioeconomy, green chemistry, high resource and ecoefficiency as well as high occupational health and safety fostered the development of © XXXX American Chemical Society

non-isocyanate polyurethanes. Among the great variety of synthetic routes to non-isocyanate polyurethanes reported in the literature, the formation of polyhydroxyurethanes (PHUs) via the ring-opening aminolysis of polyfunctional cyclic carbonates with polyamines is attractive, since cyclic carbonates are readily available by chemical fixation of the greenhouse gas carbon dioxide with a great variety of synthetic and biobased epoxy compounds. As compared to conventional polyurethanes, PHUs do not contain thermally labile allophanate or biuret groups. Moreover, the presence of hydroxyl groups within the repeating unit accounts for enhanced hydrophilicity and adhesion via hydrogen bonding. Hence, PHUs exhibit improved thermal stability, higher Received: September 4, 2018 Revised: November 8, 2018

A

DOI: 10.1021/acs.macromol.8b01908 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

aliphatic diamines such as 1,12-diaminododecane afforded a semicrystalline PHU melting at 141 °C.24 While all commercially available thermoplastic polyurethane elastomers are derived from isocyanates, the PHU chemistry allows for an isocyanate-free synthesis route and the formation of PHU, which is not feasible in conventional PUR synthesis owing to the extremely rapid reaction of isocyanate and hydroxyl groups. The first thermoplastic polyhydroxyurethane elastomers (TPHEs) were prepared by Leitsch and Beniah who achieved nanophase-separation of hard and soft PHU segments. Typically, the molar masses varied between 3 and 10 kg mol−1. These TPHEs were completely amorphous, and tensile testing was performed on test specimens cut out from sheets prepared by hot compaction in a press. The ratio of hard and soft PHU segments governed the Young’s moduli, which varied between 2.5 and 230 MPa, while the elongation at break ranged from 10 to well above 2000%.25,26 The authors stated that nanophase-separation, which is crucial for tailoring thermoplastic elastomers, may be limited in thermoplastic polyhydroxyurethane elastomers containing polyether-based PHU soft segments because of strong hydrogen bonding between hard and soft segments.25 In a more recent publication, Beniah et al. reported on a segmented PHU containing amide groups within the hard segments, which markedly enhanced nanophase-separation and in turn improved the elasticity of these amorphous PHUs.27 To the best of our knowledge, no attempts have been made so far to prepare semicrystalline thermoplastic polyhydroxyurethane elastomers (TPHEs), in which the crystallization of the hard segments accounts for the phase-separation.28 Moreover, the syntheses of many PHU thermoplastics reported in the literature require the use of solvents in conjunction with rather long reaction times exceeding 3 h as well as laborious monomer purification by chromatography. In view of both monomer scale-up and industrially feasible TPHE synthesis, it is highly attractive to develop solvent-free melt-phase polyaddition reaction with much shorter reaction times typical for reactive extrusion processes. Herein we report on tailoring semicrystalline PHUs via the solvent-free polyaddition of biobased difunctional cyclic carbonates with 1,12-diaminododecane using a twin-screw compounder. Particularly, we examine by means of calorimetric measurements in combination with small-angle X-ray scattering and electron microscopy how the molecular architectures of semicrystalline PHU and the processing conditions influence thermal, morphological, and mechanical properties of injection-molded PHU in order to gain better understanding and control of PHU crystallization. On the basis of this insight, the ratio of hard and soft segments is balanced to create the first semicrystalline PHU thermoplastic elastomers (TPHE) containing crystallizable PHU hard segments aiming at achieving property profiles competitive with those of conventional isocyanate-based thermoplastic polyurethane elastomers (TPEs). Moreover, filaments of tailored semicrystalline PHUs are employed in extrusionbased additive manufacturing by means of fused filament fabrication or fused deposition modeling to 3D-print PHU parts.

resistance to nonpolar chemical solvents, and increased adhesion as well as wear resistance.6,8−11 Additionally, the presence of hydroxyl groups in the PHU backbone enables facile polymer-analogous functionalization and grafting of other polymers. Although cyclic carbonates with larger ring sizes exhibit considerably higher reactivity,12,13 five-membered cyclic carbonates are widely applied in PHU synthesis, since they are readily available even in commercial scale via the reaction of carbon dioxide with either polyfunctional polyols containing 1,2-diol units or epoxides.8,14 It is generally known that high purity of the starting materials and ensuring precise stoichiometry are crucial for achieving high molar masses in step-growth polymerization of difunctional monomers. Since pure diamines are widely available, many studies on the formation of PHU thermoplastics have focused on the polymerization of biobased cyclic dicarbonates.15−21 The synthesis of the first linear PHUs dates back to pioneering work by Endo and co-workers in 1993. They employed difunctional cyclic carbonates derived from diepoxides to prepare PHU with molar masses varying between 20 and 30 kg mol−1 by polyaddition with diamines in dimethylacetamide as solvent at 100 °C for the duration of 24 h.22 Benyahya prepared PHUs from terephthaloyl dicyclocarbonate with various diamines (DMF, 75 °C, 48 h) to determine structure−property relationships by analyzing the resulting glass transition temperatures varying from 4 to 78 °C as a function of the PHU molecular architectures.23 Carré prepared cyclic dicarbonates by esterification of biobased dicarboxylic acid with the glycerol-derived monocarbonate, which yielded fully biobased PHUs in a solvent-free polyaddition process (75 °C, 2 h) producing PHUs with molar masses ranging from 6 to 9 kg mol−1.17 According to Lamarzelle, an electron-withdrawing adjacent ester group drastically enhanced the reactivity of a cyclic carbonate, accounting for the absence of amidation as side-reaction when choosing the appropriate reaction conditions (DMF, 70 °C, 7 days) and varying the reaction temperature between 25 and 70 °C.19 Similarly, Duval obtained a dicarbonate monomer by esterification of sebacic acid with glycerol carbonate in the presence of DCC/DMAP. Typically, the solventless melt polyaddition with various diamines was achieved at relatively high temperatures of 150 °C for the duration of 16 h in the presence of 1,5,7-triazabicyclo[4.4.0]dec-5-ene (TBD) as catalyst. However, most of the resulting PHUs were insoluble, thereby drastically impairing their molar mass determination. Some of these PHUs were reported to be semicrystalline with melting temperatures varying between 102 and 165 °C.18 In another strategy, Maisonneuve prepared fatty-acid-based dicarbonates from dimerized methyl-10-undecenoate, which was epoxidized and carbonated. They prepared various PHUs (bulk, 70 to 140 °C, 1−13 days) with molar masses up to 31 kg mol−1, some of which exhibited broad melting ranges around 115 °C.21 It is noteworthy that in the literature the reports on amorphous PHUs outnumber by far the reports on semicrystalline PHU. Moreover, in the cases of the last two examples, crystallization is strongly favored by the presence of ester and amide groups within the molecular framework of cyclic carbonate monomer. In a recent advance, Schmidt et al. prepared PHUs by polyaddition of various aliphatic diamines with butadiene dicarbonate, which represents the smallest difunctional cyclic carbonate feasible. While polyaddition with 1,6-diaminohexane and 1,8-diaminooctane yielded amorphous products, the corresponding polyaddition with longer chain

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EXPERIMENTAL SECTION

Materials. 1,4-Butanediol diglycidyl ether (1a, ERISYS GE-21), 1,6-hexanediol diglycidyl ether (1b, ERISYS GE-25), and neopentyldiol diglycidyl ether (1c, ERISYS GE-20) were kindly provided B


Article

Macromolecules by CVC Emerald Materials. Resorcinoldiol diglycidyl ether (1d), PEG diglycidyl ether (1g, Mn = 500 g mol−1), allyl glycidyl ether (99%), 1,5,7-triazabicyclo[4.4.0]dec-5-ene (TBD, 98%), 1,12-diaminododecane (DDA, 98%), 4,9-dioxa-1,12-dodecanediamine (DODA, 99%), tetrabutylammonium bromide (TBAB, 99%), and 3-chloroperbenzoic acid (mCPBA, 77%) were purchased from Sigma-Aldrich. 1,6Hexamethylenediamine (HMDA, 98%) and 1,9-decadiene were obtained from SAF and abcr, respectively. Carbon dioxide (N45) was purchased from Air Liquide. Transparent polylactide (PLA) filament with 2.85 mm thickness was obtained from Ultimaker. Characterization. NMR spectra were recorded in deuterated chloroform or DMSO on an ARX 300 spectrometer from Bruker at room temperature. The chemical shifts were referenced to the solvent signals. DSC measurements were performed using a PerkinElmer’s Pyris 1 with a heating and cooling rate of 10 K min−1, if not otherwise indicated. Atomic force microscopy was carried out at room temperature on a Multimode AFM IIIa from VEECO in tapping mode, using injection-molded samples after the preparation of a flat surface at −120 °C using a cryomicrotome ULTRACUT UCT from Leica. Small-angle X-ray scattering (SAXS) experiments were carried out with a Kratky camera attached to a conventional Cu Kα X-ray source employing a position sensitive metal wire detector and a temperature controlled sample holder. After background correction, a deconvolution algorithm was applied on the slit-smeared scattering data.29 DMA measurements were performed on a Q800 from TA Instruments using a dual cantilever measuring system. After equilibration at −40 °C for 5 min, the rectangular sample (80 × 12 × 3 mm) was heated with a rate of 3 K min−1, a deformation of 0.1%, and a frequency of 1 Hz. The hysteresis measurements were also conducted on this Q800 with a film tension clamp setup. The injection-molded film was strained at 21 °C to 100% with a strain rate of 50% min−1 and then relaxed to 0% elongation at 3.33% min−1. The hysteresis was measured for five consecutive cycles. Tensile testing was performed on a Zwick Z005 (Ulm, Germany, ISO 527-1/2) with a strain rate of 50 mm min−1. The mechanical properties such as elastic modulus, tensile strength, and breaking elongation at room temperature were determined by taking the statistical average of four to six test specimens (5A). SEC measurements were performed in N,N-dimethylacetamide (DMAc) on a PSS Agilent 1200 Series with connected refractive index detector. For separation at room temperatur,e SDV columns (5 μm; 100, 1000, and 10 000 Å) were used, and the calibration was carried out with poly(methyl methacrylate). PerkinElmer’s Pyris TGA 4000 was used for the thermogravimetric analysis (TGA). The sample was heated under air from 50 °C up to 650 °C with a heating rate of 10 K min−1. Scanning electron microscopy (SEM) pictures were taken on an Amray 1810 using an electron acceleration of 15−20 kV at a working distance of 10 mm. To prevent electrostatic charging of the samples, a small coating of gold (10−20 nm) was applied. Purification of Technical Grade Glycidyl Ethers. Technical grades of 1a, 1b, and 1c are typically obtained by glycidylization of the respective diols with epichlorhydrin. To remove high molar mass byproducts resulting from advancement reactions, the diepoxides were purified by distilliation (Figure S1). The respective conditions, yields, and resulting purities are listed in Table 1. Purities for the technical grades before the distillation cannot be determined using the 1 H NMR spectra because of overlapping signals. As determined from the 1H NMR spectra, the products contain 21−28 mol % monogylcidyl ethers after distillation, which have similar boiling points like the diglycidyl ethers. The monofunctional components will

eventually be removed through precipitation of the dicarbonate in the subsequent synthetic step. Synthesis of 2,2′-[(Oxybis(methylene)]bis(oxirane) (1e). Allyl glycidyl ether (42.93 g, 376 mmol) was dissolved in methylene chloride (50 mL) and cooled with an ice bath. A solution of metachloroperbenzoic acid, mCPBA (64.69 g, 289 mmol) in methylene chloride (300 mL), was slowly added over a period of 1 h using a dropping funnel. Then, the ice bath was removed, and more mCPBA (32.67 g, 146 mmol) was added. The reaction mixture was stirred for another 24 h before the precipitates were filtered off and a solution of sodium thiosulfate (66 mL saturated Na2S2O3 solution diluted with 134 mL distilled water) was added. After the aqueous phase was vigorously stirred for 15 min, it was separated, and the organic phase was washed with aqueous NaOH (0.5 M, 200 mL). The organic phase was dried over Na2SO4, and the solvent was stripped off under reduced pressure to yield 1e as a colorless liquid (1e, 23.66 g, 182 mmol, 48%). Synthesis of 1,6-Di(oxiran-2-yl)hexane (1f). 1,9-Decadiene (50.70 g, 367 mmol) was dissolved in methylene chloride (600 mL) and cooled with an ice bath prior to the addition of solid mCPBA (196.19 g, 1137 mmol) in two portions. Half of the stoichiometric amount was added, and an hour later, the ice bath was removed. Then, the second half portion was added at room temperature, and the reaction was stirred for another 24 h before the resulting precipitate of m-chlorobenzoic acid was filtered off. The organic phase was then washed twice with sodium thiosulfate solution (half saturated, 400 mL) and twice with sodium hydroxide (0.5 M, 400 mL) and dried over NaSO4. The solvent was stripped off under reduced pressure to yield 1f as a colorless liquid (57.34 g, 372 mmol, 92%). General Procedure for the Carbonation of Diepoxides. The diepoxide was fed to a stainless-steel autoclave (MiniClave, 200 mL or BEP-280, type 4, 2 L from Büchi Glas Uster, Switzerland). After 1−2 wt % tetrabutylammonium bromide (TBAB) was added as catalysts and the reactor was flushed three times with 30 bar CO2, the autoclave was pressurized with 30 bar CO2 as soon as the reaction temperature was reached. The preferred stirring speed was 350 rpm. Except for 2g, all products were precipitated from solution and yielded crystalline grades with high purity as determined from the 1H NMR spectra. All relevant reaction conditions and precipitation parameters are listed in Table 2. The thermal properties as determined by differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA) are summarized in Table 3. The corresponding 1H/13C NMR spectra are displayed in Figures S3−23). Polymerization Procedures. Polyaddition was performed in bulk using either a twin-screw mini compounder (Process 1) or a stirred glass flask (Process 2). 1. Polyaddition in a Twin-Screw Mini Compounder. Cyclic dicarbonates and diamines (and 1,5,7-triazabicyclo[4.4.0]dec-5-ene in the case of 3a_cat) were mixed in a beaker at 120 °C for 10 min and then poured onto a Teflon plate. After solidification by crystallization, the resulting solid was crushed into pieces and fed into the preheated Micro 15 cm3 Twin Screw Compounder from XPLORE (DSM, Geleen, Netherlands). The viscous mixtures of 3d and 3f were fed directly in to the compounder, as they do not crystallize. The screw rotation speed was 120 rpm. The polymerization conditions are summarized in Table 4. The resulting polymers were subsequently injection-molded (9 bar, 10 s) to produce a testing specimen for the mechanical analysis. The melt temperature was identical to the final polymerization temperature, while the mold was usually kept at room temperature, except for 3e with a mold temperature of 40 °C. 2. Polymerization in a Stirred Glass Flask. The respective monomers were added to a flask and stirred at elevated temperatures under a nitrogen atmosphere using a mechanical stirrer. The polymerization runs are listed in Table 4. Filament Fabrication and 3D Printing. The PHU 3e was fed into the twin-screw compounder, and after 5 min at 130 °C and 120 rpm, the nozzle was opened, and the extruded polymer filament was wound up on a spool. Because of the fast crystallization of PHU 3e, the strand solidified rapidly without further cooling. By this technique,

Table 1. Distillation of Technical Grade Glycidyl Ethers 1a 1b 1c

m [g]

p [mbar]

T [°C]

yield [g]

puritya [%]

1015.6 1017.7 992.9

0.8 0.4 0.5

122 140 115

459.4 497.4 618.9

79 72 72

a

Purity after distillation as determined from 1H NMR. C


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Macromolecules Table 2. Carbonation of the Diepoxides (2a−g) T

t

product

[g]

[mmol]

[g]

[mmol]

[°C]

[h]

precipitation conditions

[g]

2a 2b 2c 2d 2e 2f 2g

457.92 389.1 522.6 355.9 17.2 97.6 340.5

1901 1307 1879 1601 132 574 681

9.2 7.8 10.4 7.12 0.3 1.9 3.4

29 24 32.3 33.2 0.9 5.9 6.8

130 130 130 130; 140 120 120 120

17 16.5 16 1; 14.5 18 22 36

450 mL acetone, 7 days at 7 °C 450 mL acetone, 24 h at −20 °C 660 mL MeOH, 20 h at r.t. 400 mL acetone, 20 h at −20 °C 50 mL acetone, 40 mL Et2O, 48 h at −20 °C 300 mL acetone, 400 mL Et2O, 48 h at −20 °C -

288.8 (52%a) 159.5 (38%1) 235.6 (41%1) 375.8 (76%) 20.3 (70%) 83.9 (57%) 131.7 (98%)

diepoxide

TBAB

a

yield

The respective amounts of monoglycidyl ether in the starting material were not taken into account.

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RESULTS AND DISCUSSION Solventless Tailoring of Thermoplastic PHUs. As illustrated in Scheme 1, various diglycidylethers 1 were converted into the corresponding difunctional cyclic carbonates 2 by chemical carbon dioxide fixation followed by meltphase polyaddition with 1,12-diaminododecane to prepare semicrystalline polyhydroxyurethane (PHU) thermoplastics 3. The biobased diglycidyl ethers 1a−d were readily gained by glycidylization,30 preferably using glycerol-derived epichlorhydrin31 and biobased diols.32−35 To ensure high monomer purity, the higher molar mass byproducts, which are present in technical grade epoxy resins, were removed by vacuum distillation of 1a−c (Figure S1). The carbonation of the epoxides was performed in a stainless-steel reactor by means of the chemical carbon dioxide fixation in the presence of the tetrabutyl ammonium bromide (TBAB) as catalyst under carbon dioxide pressure (30 bar) as described in the literature.36 The biobased cyclic carbonate 4,4′-(oxybis(methylene))bis(1,3-dioxolan-2-one) (2e), also available by transesterification of diglycerol with dimethyl carbonate,15,16 was prepared by epoxidation of the allyl glycidyl ether 1e and subsequent carbonation. Decarboxylation of ricinoleic-acidderived undecylenic acid affords biobased 1,9-decadiene,37−39 which after epoxidation and carbonation, yielded 4,4′-(hexane1,6-diyl)bis(1,3-dioxolan-2-one) (2f). To the best of our knowledge, 2f was never used before employment in PHU synthesis. The flexible dicarbonate-terminated oligo(ethylene oxide) 2g was derived by carbon dioxide fixation of diglycidylether-terminated oligo(ethylene oxide) with a number-average molar mass Mn = 500 g mol−1. The facile purification by recrystallization enabled easy scaleup of the 2a−f preparation (up to 375 g scale), meeting the

Table 3. Thermal Properties of Cyclic Carbonate Monomers dicarbonate

Tma [°C]

ΔHma [J g−1]

Tdb [°C]

2a 2b 2c 2d 2e 2f

80 67 90 60−115 74 40−75

99 121 110 55 101 124

281 314 302 270 271 333

a

Melting temperature and melting enthalpy as determined by DSC (−50−120 °C, 10 K min−1, first heating cycle). bDecomposition temperature as determined by TGA (50−650 °C, 10 K min−1, air).

filaments were obtained with diameters of 2.2 ± 0.6 mm and a length of up to 40 cm. This procedure was repeated three times, and the ends of resulting filaments were joined together by heating to obtain a single filament with a total length of about 1.50 m. The 3D printing of the PHU 3e filament was carried out using an Ultimaker 2+ FDM printer from Ultimaker (Geldermalsen, Netherlands). The 3D Object was printed on a glass build platform without any surface modification or addition of adhesion promoters. The build platform temperature was 50 °C, and the nozzle temperature was 190 °C. The diameter of the nozzle was 0.8 mm. The g-code generation was carried out using the software Ultimaker Cura 2.3. Herein, the wall thickness was 1.0 mm with an infill of 20% and a layer height of 0.15 mm. The print speed of 50% was kept constant, while the flow rate varied between 96 and 110% because of variations of the filament diameter to receive a constant material flow. In case of the PLA 3D printing as benchmark, the build plate and the nozzle temperature were set to 60 and 210 °C, respectively, and the material flow rate was held constant at 100%, while the other parameters remained the same.

Table 4. PHU Syntheses by Melt-Phase Polyaddition product

procedure

m (dicarbonate) [g]

m (diamine) [g]

T [°C]

t [min]

3a 3a_cata 3b 3c 3d 3e 3f 2a+DODDA 2a+HMDA 3a_kpg prepol prepol-co-3e

(1) (1) (1) (1) (1) (1) (1) (2) (2) (2) (2) (1)

11.389 (2a) 10.658 (2a) 12.318 (2b) 11.464 (2c) 11.475 (2d) 10.882 (2e) 10.93097 (2f) 29.35 (2a) 7.111 (2a) 10.21 (2a) 29.2560 (2g) 3.26862 (2e)

8.022 (DDA) 7.357 (DDA) 7.912 (DDA) 7.702 (DDA) 7.562 (DDA) 9.989 (DDA) 8.47990 (DDA) 20.66 (DODDA) 2.847 (HMDA) 7.19 (DDA) 11.6580 (DODDA) 13.34643 (prepol); 2.45783 (DDA)

100; 110 100; 110 100; 110 100; 110 120; 130 110; 120; 135; 140 100; 110; 120; 130 100; 110 90; 110 100; 110; 120; 130 100 100; 110; 120; 125; 130

45; 20 68; 27 45; 20 70; 30 8; 10 5; 20; 7; 7 15; 15; 15; 85 30; 900 420; 30 90; 180; 120; 300 90 15; 15; 25; 10; 65

a

Contains 0.184 g of 1,5,7-triazabicyclo[4.4.0]dec-5-ene (TBD, 1 wt %). D


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Macromolecules

Scheme 1. Reaction Pathway toward Polyhydroxyurethanes (PHU) 3 by Polyaddition of 1,12-Diaminododecane with Various Difunctional Cyclic Carbonates 2 Derived from Diepoxides 1a

a

(a) Reaction conditions for cyclic carbonate formation: 120−140°C, 1−2 wt.% tetrabutyl ammonium bromide (TBAB) as catalyst, 30 bar CO2, 16−26 h; (b) reaction conditions of melt-phase polyaddition in a twin-screw compounder: nitrogen atmosphere, 100−130°C, 18−130 min.

Table 5. Polyaddition Reaction Conditions and PHU Properties As Determined by SEC, DSC, and TGA Measurements PHU

trcta [min]

Trctb [°C]

Mnc [kg mol−1]

Mwc [kg mol−1]

Đc

Tgd [°C]

Tmd [°C]

ΔHmd [J/g]

Tdf [°C]

3a 3a_kpg 3a_cat 3b 3c 3d 3e 3f 3f_anng prepolh prepol-co-3eh

65 690 92 65 100 18 39 115

100−110 100−130 100−110 100−110 100−110 120−130 110−140 100−130

26.7 25.7 12.6 32.9 20.5 18.6 20.7 13.9

53.9 52.0 27.0 70.8 41.1 61.1 40.4 27.4

2.0 2.0 2.2 2.1 2.0 3.3 2.0 2.0

90 130

100 100−130

6.6 14.0

19.0 74.2

2.9 5.3

2 n.m. n.m. 4 10 35 34e 20e 23 −32 −30

40−80 n.m. n.m. 40−80 40−60 90−120 75−90 60−90

40 n.m. n.m. 40 39 55 18 10

294 n.m. 200/144 304 298 298 285 284 n.m. n.m. 286

Reaction time. bReaction temperature (precise time/temperature program given in Table 4 and in Figure S2). cSEC (DMAc, 85 °C, PMMA standard). dDSC (−50−120 °C, 10 K min−1, first heating cycle). eSecond heating cycle. fTGA (50−650 °C, 10 K min−1, air). gSpecimen of PHU 3f after annealing (60 °C, 72 h). hThe prepolymer route will be discussed in the section “Semicrystalline Thermoplastic PHU Elastomers”; n.m. = not measured. a

stringent high-purity demand imperative for step-growth polymerizations. High monomer purities were paralleled by higher melting temperatures as compared to those reported in the literature (Figure S51). For instance, 1a showed a melting peak at 80 °C, which was more than 19 °C greater as compared to the literature.22 Moreover, 2e melted at 74 °C, which was 8−9 °C above the literature values.15,16 While 2b and 2c were known as yellowish or clear liquid, respectively,40,41 both compounds were obtained as colorless crystalline solids. The high monomer purities were also confirmed by 1H and 13C NMR spectroscopy (Figures S11− 22). The melt-phase polyaddition of 2a−f with 1,12diaminnododecane was performed by means of a twin-screw compounder. Besides ecological advantages of this solvent-free synthesis, this method allowed for qualitative in situ conversion monitoring using the implemented force control unit, since the measured force is directly linked to the melt viscosity of the reaction mixture. The resulting PHUs 3a−f were analyzed by means of 1H NMR spectroscopy using signal assignments reported previously (Figures S24−35).15,24,41,42 All PHU molar masses were determined by means of SEC (DMAc, PMMA standard), which showed monomodal distributions (see Figures S36−47) and number-average molar masses varying between 13.6 and 32.9 kg mol−1 (see Table 5). As expected for step-growth polymerization, the polydispersities (Đ) of all PHU homopolymers were around 2 with exception of the higher polydispersity of 3d. This may be attributed to side-reactions occurring at high shear stress, which were present in the case of 3d as a consequence of the rigid, aromatic building block 2d. The PHU 3a was chosen to

investigate the influence of polymerization conditions on PHU molar mass and thermal behavior. Using the same starting materials, 3a_kpg was prepared by mechanical stirring. Neither increased reaction temperature nor 10-fold increased reaction time significantly affected PHU molar masses (see Table 5), although the polyaddition reaction was incomplete, as evidenced by the presence of the remaining functional groups (Figures S24 and S30). This can be attributed to decreasing reaction rate with decreasing concentrations of both functional groups and also to higher viscosity resulting from increased hydrogen bonding limiting the advancement reaction as proposed by Blain et al.43 In 3a_cat, 1 wt % 1,5,7-triazabicyclo[4.4.0]dec-5-ene (TBD), a known catalyst for the aminolysis of cyclic carbonates,24,44,45 was added using the reaction conditions identical to those of the 3a preparation. The TBD addition drastically impaired PHU molar mass, which was lowered from 26.7 down to 12.6 kg mol−1. This is in contrast to reports by Chen et al. stating that TBD catalyst addition at room temperature markedly increased PHU molar mass.46 However, to the best of our knowledge, no such study was conducted at elevated temperatures typical for this melt-phase polyaddition. This drastic PHU molar mass loss encountered at polyaddition temperatures of 100−110 °C in the presence of 1 wt % TBD was attributed to the catalysis of side-reactions, which become more significant at elevated temperature. This was paralleled by the much lower decomposition temperature of 3a_cat at 200 °C as compared to 294 °C of the catalyst-free 3a. Moreover, 3a_cat showed a lower weight loss of only 1.2% at 144 °C in thermogravimetric analysis (Figure S49). Various E


Article

Macromolecules Table 6. Mechanical Properties of Injection-Molded PHUs according to ISO 527-1/2_5A, 50 mm min−1 PHU

Young’s modulus [MPa]

3a 3a_cat 3b 3c 3d 3e 3f 3f_ann

220 ± 30 240 ± 10 225 ± 18 378 ± 16 1780 ± 120 1430 ± 130 4±2 220 ± 30

σya [MPa] ± ± ± ± 46 ± -

11.4 11.1 10.4 15.6

εyb [%] ± ± ± ± 3.8 ± -

0.7 0.4 0.9 0.9

7.9 7.40 7.0 4.83

5

0.3 0.11 0.4 0.13 0.3

σbc [MPa]

εbd [%]

17 ± 2 10.6 ± 0.2 18 ± 2 8±5 9.0 ± 0.4 28 ± 9 27 ± 5 31 ± 4

300 ± 40 9.1 ± 0.6 340 ± 40 25 ± 16 0.36 ± 0.01 120 ± 60 470 ± 70 180 ± 30

a

Yield stress. bElongation at yield. cTensile strength at break. dElongation at break.

Figure 1. (a) DSC thermograms of injection-molded PHU 3a at various heating rates (−50−120 °C, first heating cycle). (b) DSC thermograms (−50−120 °C, 10 K min−1) of samples from PHU 3a after being cooled from 120 to 21 °C at 10 K min−1 and subsequently stored at room temperature or at elevated temperature in on oven. The samples were measured after 24 h via DSC (−50−120 °C, 10 K min−1, first heating cycle). (c) Correlation functions as calculated from SAXS intensities for PHU 3a measured at 25 and 70 °C. The elevated temperature was achieved through a stepwise heating with an effective heating rate of 0.2 K min−1. The crystallite thicknesses dc were determined from the correlation function as the z-value of the intersection point between the straight part of the initial slope and the baseline at the first minimum.50

straining (Figure S55). This stress-induced crystallization was observed above 180% elongation and was paralleled by significantly increased tensile strength (Figure S58). This strain-hardening by strain-induced crystallization was observed for 3a, 3b, and also for 3f, which was amorphous prior to straining. Moreover, the crystallization of 3f was also induced by annealing at elevated temperature. One set of tensile test bars of 3f was annealed at 60 °C for 72 h. From Table 5, it is apparent that annealing rendered the corresponding 3f_ann semicrystalline (Figure S52), paralleled by a significantly improved Young’s modulus (4 to 220 MPa) and a higher tensile strength (27 to 31 MPa) at the expense of elongation at break, which was reduced from 470 to 180% (see Table 6). Crystallinity of PHU Thermoplastics. DSC measurements of semicrystalline PHUs were performed at heating and cooling rates of 10 K min−1. Under these conditions, only 3e crystallized during cooling with a crystallization temperature of 67 °C, while 3a, 3b, and 3c crystallized slowly over prolonged periods of time. In the case of 3a, samples were cooled from 120 to 21 °C at 10 K min−1 and stored for 10 and 30 min as well as 1, 4, 24, and 48 h before measuring the next heating cycle. The resulting melting enthalpies were determined to 0, 23, 30, 33, 36, and 39 J g−1, respectively, showing that the major part of this “slow crystallization” at room temperature takes place within the first hour. As shown in Figure 1a for PHU 3a, these semicrystalline PHUs usually exhibited multiple maxima in their DSC thermograms when measured at 10 K min−1. This was attributed to recrystallization occurring during

publications have already elaborated on side-reactions in PHU synthesis as a limiting factor for achieving high molar masses.19,47−49 In the light of these results, it can be concluded that catalysts such as TBD are less suited for melt-phase bulk polyadditions at elevated temperatures. From Table 6, it is apparent that this significant PHU molar mass loss drastically impaired mechanical properties. As compared to 3a_cat, the sample 3a prepared in the absence of catalyst exhibited a 30fold higher elongation at break εb. Therefore, all other meltphase reactions were performed in the absence of catalysts. After injection molding at the respective reaction temperature (mold: room temperature and 40 °C in the case of 3e) and subsequent cooling to room temperature, the PHUs 3a, 3b, 3c, and 3e were semicrystalline, whereas 3d and 3f were amorphous with glass temperatures around room temperature, as determined by DSC (see Table 5). Accordingly, 3d with a glass transition temperature above room temperature was stiff but very brittle with an εb below 1%, whereas 3f with a glass transition temperature below room temperature was rather flexible, exhibiting a high εb and low Young’s modulus (see Table 6). The PHUs 3a, 3b, 3c, and 3e exhibited a yield point as expected for semicrystalline thermoplastics. Among these four samples, 3e showed highest stiffness as reflected by its Young’s modulus around 1430 ± 130 MPa and the highest yield stress of 46 ± 5 MPa. In the case of 3a, the thermal analysis revealed that the strained section of the tested tensile bars showed a much higher crystallization enthalpy of 57 J g−1 as compared to 40 J g−1 for the original sample prior to F


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which in turn allows for recrystallization of these polymer chains. During this step, larger crystallite thicknesses must emerge, which is realized by the incorporation of the “cyclic carbonate part” into the crystallites. When replacing 1,12-dodecanediamine (DDA) with 4,9dioxa-1,12-dodecanediamine (DODA), the polyaddition with 2a produced the completely amorphous PHU 2a+DODA (Mn = 29.2 g mol−1, Đ = 2.1, Tg = −20). Despite the similar length of the amine building block, the two additional ether groups completely prevented polymer crystallization, which confirms that the choice of a well crystallizable amine component is crucial. Likewise, substituting DDA for 1,6-hexamethylenediamine (HMDA) rendered the resulting PHU 2a+HMDA (Mn = 26.9 g mol−1, D̵ = 2.2, Tg = −5) amorphous, since the ability to crystallize was much lower for the shorter aliphatic diamine. The corresponding structures of amorphous and semicrystalline PHUs are displayed in Figure 3.

the heating process, as is apparent from the DSC traces displayed in Figure 1a measured at a slower (1 K min−1) and a faster heating rate (30 K min−1). At a heating rate of 30 K min−1, only one melting peak was detected, since there was not enough time for recrystallization, while the slow heating rate afforded multiple melting peaks and a local minimum, which also go to heat flow values below the baseline and therefore must represent recrystallization. In fact, the thermal profiles can be adjusted to some extent via controlled annealing of the respective PHU samples as shown in Figure 1b for PHU 3a. Higher annealing temperatures afforded improved chain mobility and therefore allowed for the formation of larger crystallites, which melted at higher temperatures according to the Gibbs−Thomson equation, while smaller crystallites became increasingly unstable. PHU 3a annealing at 80 °C accounted for the highest melting temperature ranges, and at higher annealing temperature, the crystallization ceased completely. As illustrated in Figure 1c, small-angle X-ray scattering (SAXS) measurements at 25 and 70 °C revealed crystallite thicknesses dc of 2.4 and 4.4 nm, respectively, thereby confirming in accord with the Gibbs−Thomson relationship that the higher temperature afforded larger crystallite thicknesses. Furthermore, it can be seen in Figure 2 that the smaller crystallite thickness of 2.4 nm matches the length of “the amine

Figure 3. Repeat units of PHUs: 3a (above), 2a+DODDA (center), and 2a+HMDA (below).

Clearly, the presence of regioirregularities in linear PHUs markedly impaired PHU crystallization. As a rule, the impact of regioirregularities is more pronounced when using smaller difunctional cyclic carbonate monomers. In accord with earlier observations, small carbonate monomers only produced semicrystalline PHU by polyaddition with long-chain aliphatic diamines such as DDA and by incorporation of amide groups in the PHU backbone.24 With increasing chain length of the difunctional cyclic carbonate, however, the impact of regioirregularities significantly declined. Hence, long-chain symmetric difunctional cyclic carbonate monomers containing long-chain oligomethylene sequences or amide or ester groups, respectively, also promote formation of semicrystalline PHU.21,45 Semicrystalline Thermoplastic PHU Elastomers (TPHE). According to Torkelson and co-workers, who investigated amorphous polyether-based TPHE, strong hydrogen bonding of hard-segment hydroxyl groups to the softsegment ether oxygen atoms significantly impaired nanophaseseparation, which is the key to tunable mechanical properties and high resilience. In fact, single-phase polyether-based PHU containing oligo(ethylene oxide) segments were not elastomeric but flowed even under the force of gravity. On the contrary, nanophase-separated oligo(tetramethylene oxide)based TPHEs were tailored as broad-temperature-range acoustic and vibration damping materials unparalleled by conventional TPUs.25 In another strategy toward nanophaseseparated TPHE, inspired by isocyanate-based TPUs containing semicrystalline hard segments,28 the semicrystalline PHU

Figure 2. Two constitutional isomers of the 3a repeat units with two secondary (above) and two primary hydroxyl groups (below), respectively, in all-trans conformation. The third possible isomer with one primary and one secondary hydroxyl group is not shown.

part” of the repeat unit (Luau = 2.3 nm), while the larger crystallite thickness of 4.4 nm matches the length of the full repeat unit including the “cyclic carbonate part” (Lmax = 4.5 nm). The presence of constitutional isomeric repeat units, accounted for by the poor regioselectivity of the cyclic carbonate aminolysis in accordance with the two feasible pathways for the ring-opening reaction, causes irregularities along the polymer backbone, which are located in the “cyclic carbonate part” of the repeat unit and hamper crystallization. The ratio between secondary and primary hydroxyl groups for all polymers within this PHU family was roughly 2:1, as the hydroxyurethane with a secondary hydroxyl group is known to be more stable and therefore typically represents the major product.51 On the basis of the results from the SAXS measurements, it can be concluded that at lower crystallization temperatures, “the amine part” predominately crystallizes without including the “cyclic carbonate part” into the crystallites, thereby causing low crystallite thicknesses paralleled by low melting temperatures. At higher temperatures, these small crystallites naturally become unstable and melt, G


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Macromolecules 3e, melting at 90−120 °C, was incorporated as hard segment into a flexible PHU backbone. Figure 4 displays the two-step

second stage was performed by means of a twin-screw melt compounder. The resulting TPHE prepol-co-3e exhibited a much higher molar mass (Mn = 14.0 g mol−1, Mw = 74.3 g mol−1, Đ = 5.3) while preserving a monomodal molar mass distribution (Figures S46 and 47). As shown in Table 5, the molar mass of prepol-co-3e turns out slightly lower when compared to the homopolymers. This circumstance is probably caused by incorporation of the flexible PEG-based dicarbonate 2g, which typically exhibits a functionality slightly lower than 2. The glass temperatures of prepol (−32 °C) and prepol-co3e (−30 °C) are very similar and unaffected by the incorporation of the semicrystalline PHU 3e having a much higher glass temperature of 34 °C. This provides further experimental evidence for phase-separation. The Tg of the hard phase was not detected by DSC measurements, which revealed a broad melting range between 60 and 90 °C (ΔHm = 10 J g−1). Hence, the crystallites formed by crystallization of the hard segments served as thermally reversible cross-links. From the dynamic-mechanic analysis displayed in Figure 5a, it became apparent that prepol-co-3e maintained its dimensional stability up to 60 °C. Unlike the homopolymer 3e, the segmented PHU shows a shoulder in the SAXS trace (see Figure 5b) at around s = 0.06 nm−1. This again indicates nanophase-separation with domain sizes around 17 nm. Furthermore, as is apparent from Figure 5c, this nanophaseseparation of prepol-co-3e was confirmed by means of atomic force microscopy (AFM). Tensile tests on a rectangular film specimen were conducted to examine the elastic properties of prepol-co-3e. In cyclic loading, the sample was strained to 100% for five consecutive cycles. As is apparent from Figure 6, the strain recovery by mechanical hysteresis for the first cycle was calculated according to eq 1 to 74% and then drastically decreased to 32, 30, 29, and 28% in the four subsequent cycles. During the “unloading”, the measured stress became negative for deformations lower than 33% because of a bulge formation of the sample. This bulge formation and the large hysteresis in the first cycle indicate slippage of polymer chains in accordance with reports in the literature for other TPEs.25,27,52−54

Figure 4. Two-step synthesis of the nanophase-separated semicrystalline TPHE (prepol-co-3a): In the first step, (a) diamine-terminated flexible amorphous PHU prepolymer (prepol) was formed as a soft segment by nonstoichiometric melt-phase polyaddition of 2g with 4,9dioxa-1,12-dodecanediamine (DODA) (100 °C, 1.5 h, nitrogen atmosphere, n ≈ 8) in a stirred reaction; in the second step, (b) the subsequent stoichiometric co-polyaddition of prepol with in situ formed PHU 3e hard segments was performed in a twin-screw melt compounder (100−130 °C, residence time = 130 min, nitrogen atmosphere).

polyaddition process for the preparation of TPHE prepol-co3e. Typically, in the first step, the difunctional flexible dicarbonate-terminated oligo(ethylene oxide) 2g was reacted with an excess of the flexible DODA (molar ratio of 6:7) to yield the amorphous diamine-terminated soft-segment prepol (Mn = 6.6 g mol−1, Mw = 19.0 g mol−1, Đ = 2.9). In the second step, prepol was copolymerized with stoichiometric amounts of DDA and 2e, which formed in situ 30 wt % of the semicrystalline hard segment corresponding to 3e. Taking into account much higher melt viscosities and stirring problems, the

MH =

AUC load − AUCunload *100; AUC: area under the curve AUC load (1)

Figure 5. Nanophase-separation of the TPHE prepol-co-3e; (a) DMA (−40−90 °C, 3 K min−1, 1 Hz, 1% deformation), (b) SAXS, and (c) AFM (tapping mode, phase image). H


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Figure 6. Mechanical hysteresis of TPHE prepol-co-3e; five consecutive hysteresis cycles as determined from tensile testing of injection-molded films (thickness: 0.64 mm, 21 °C, straining with 50 mm min−1 up to 100% elongation, strain release with 3.33 mm min−1 to 0%).

Figure 7. 3D-printed die made of PHU 3e using a commercial FDM printer (Ultimaker 2+). The blue color was added to highlight the printed cavities.

It is noteworthy that after the first cycle, all subsequent cycles were remarkably reproducible, and only minor “material flow” occurred. The hysteresis behavior was rather similar to that of amorphous TPHE published by Beniah et al., who incorporated an aromatic diamidodiamine monomer into the PHU backbone.27 Both examples demonstrate that TPHEs are entering the property range typical for conventional isocyanate-based TPUs,52,53 once the problem of phase intermixing by strong hydrogen bonding of hard-segment hydroxyl groups to soft-segment hydroxyl groups or oligo(ethylene oxide) segments is overcome. 3D Printing of Semicrystalline PHU Thermoplastics. Digital free-form fabrication and rapid manufacturing by means of 3D and 4D printing is attracting rapidly expanding industrial interest and expands the horizon of polymer processing with respect to customized additive manufacturing. Digital slicing of model data from a computer-designed 3D model or from 3D scans enables building in sequential layers objects of almost any shape or geometry. Today, a great variety of polymers are tailored to meet the demands of 3D printing ranging from powder fusion to stereolithography and 3D dispensing.55 Among additive manufacturing technologies, fused filament fabrication (FFF) or fused deposition modeling (FDM) is exceptionally robust with respect to additive manufacturing of thermoplastics. Typically, polymer filaments are fed into the heated extrusion print head and enable the 3D dispensing of the resulting polymer melts. The rather narrow FFF processing window requires a careful balance of processing parameters such as temperature, rheology, build speed, and path planning to achieve the desired combination of high dimensional accuracy, acceptable mechanical properties, and high surface quality. Most common commercial materials for FFF include ABS and polylactide.56−58 To date, opposite to isocyanatebased polyurethane,59 PHUs have not been employed in FFF. Among the PHU family, PHU 3e, melting at 90 to 120 °C, was chosen for FDM, since its crystallization temperature at 67 °C (DSC, 10 K min−1) ensures rapid solidification. Figure 7 displays a 3D-printed die using PHU 3e, which was 3D-printed by means of a commercial FFF printer (Ultimaker 2+) using a 3e filament, which was produced by means of melt extrusion at 130 °C.

The 3D deposition of molten polymer strands and their melt fusions requires sufficient polymer entanglement and adhesion of adjacent printed layers to prevent formation of pores and structural defects impairing mechanical properties. SEM images displayed in Figure 8 show the adjacent strands of a

Figure 8. SEM imaging of 3D-printed dice made of PHU 3e (left) and polylactide (PLA, right).

3D-printed die made of PHU 3e and polylactide (PLA). The build precision is very similar. However, at higher magnification, the gap between two strands in Figure 8 below shows that in the case of PHU 3e, strands are typically located closer together. This is attributed to strong hydrogen bonding of the hydroxyl groups in PHU 3e, which enhance adhesion between neighboring strands and thus improve interlayer adhesion and layer fusion. Another common problem encountered in FDM printing of semicrystalline polymers is warpage resulting from shrinkage due to crystallization of the deposited material. This can lead to undesirable premature separation of the 3D-printed object I


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Macromolecules from the build platform during the 3D printing process.60 Frequently, this warpage problem is solved by the incorporation of additional adhesion-promoting interlayers between the build platform and the object. In contrast, no such additional steps were necessary for 3D printing of PHU 3e. Owing to strong hydrogen bond formation of the PHU 3e, 3Dprinted PHU 3e strongly adhered to the glass build platform.61 Warpage may also be caused by a strong thermal gradient in the FDM part, which can be reduced by heating the FDM chamber or increasing thermal conductivity of the polymer through fillers.60,62 Typically, the FDM-printed objects are heated at a temperature well above the glass transition Tg to improve adhesion of adjacent strands but below the crystallization temperature Tc to ensure solidification. The low thermal transition temperatures of PHU 3e (Tg = 34 °C, Tc = 67 °C, Tm = 90−120 °C) allow for 3D printing at lower processing temperatures, which are feasible by most common 3D printers, without encountering any noticeable warpage.

to that of isocyanate-based TPE. In contrast, conventional amorphous PHU containing oligo(ethylene oxide) segments of similar chain length typically suffer from severe “phase intermixing” of soft and hard segments because of strong hydrogen bonding of hard-segment hydroxyl groups to softsegment ether groups. In the absence of nanophase-separation no mechanical hysteresis is achieved. The incorporation of semicrystalline PHU hard segments solves this miscibility problem and allows for controlled nanostructure formation. The presence of hydroxy groups is attractive with respect to 3D printing of semicrystalline PHU without encountering warpage problems. This is attributed to significantly improved adhesion between 3D-deposited layers and also between PHU and glass. Tailoring semicrystalline PHU holds great promise for 3D and 4D printing of biocompatible PHU for medical application.

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ASSOCIATED CONTENT

S Supporting Information *

CONCLUSION The solventless polyaddition of difunctional cyclic carbonates, readily derived from diepoxides by chemical fixation of the greenhouse gas carbon dioxide, with long-chain diamines such as 1,12-diaminododecane (DDA), yields a wide variety of semicrystalline polyhydroxyurethane (PHU) thermoplastics. Also, 100% biobased, semicrystalline PHU thermoplastics are feasible. Optimized and efficient purification by recrystallization affords difunctional cyclic carbonate monomers with high purity, which is imperative for achieving high molar mass in step-growth polymerization. Preferably, melt-phase polyaddition is performed in a twin-screw compounder at polymerization temperatures varied between 110 and 140 °C in the absence of catalysts, which drastically reduces PHU molar masses owing to catalysis of undesirable side-reactions, which impair the precise stoichiometry. As a function of PHU molecular architectures, PHU properties are controlled over a very wide range with PHU melting temperatures varied between 40 and 115 °C and PHU Young’s moduli ranging from 220 to 1430 MPa. According to DSC and SAXS investigations, PHU crystallization is primarily affected by the diamine chain length and annealing. The incorporation of long-chain diamines such as DDA is essential for enabling crystallization, since the aminolysis of cyclic carbonates has low regioselectity producing a mixture of primary and secondary hydroxyl groups in the PHU repeat units, which severely hamper crystallization. The “dilution” of such regioirregularities, paralleled by the increasing chain length of the aliphatic diamine, enhances PHU crystallinity. Alternatively, also the incorporation of ester and amide groups promote crystallization of PHUs. This gained insight into the PHU crystallization behavior has led to the development of a new family of PHU thermoplastic elastomers (TPHE) containing nanophase-separated soft and hard segments as evidenced by means of SAXS and AFM. The key feature is the incorporation of semicrystalline PHU hard segments into the flexible PHU backbone. In a two-step process, highly flexible diamineterminated PHU prepolymers as soft-segment precursors are prepared by nonstoichiometric polyaddition of long-chain dicarbonate-terminated oligo(ethylene oxide) with flexible diamines, followed by stoichimetric co-polyaddition with rigid short-chain dicarbonates and DDA, producing semicrystalline PHU hard segments. The resulting segmented, semicrystalline TPHE exhibited a mechanical hysteresis similar

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b01908. Compounder force unit data, 1H and 13C NMR, SEC, TGA, DSC, SAXS, DMA, and tensile testing (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]; Phone: +49 761 203 6273. ORCID

Vitalij Schimpf: 0000-0002-9232-0479 Johannes B. Max: 0000-0002-6968-1647 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge financial support of the Baden-Württemberg Stiftung GmBH (project BioMatS-02) and JONAS−Joint Research Network on Advanced Materials and Systems, BASF SE, Ludwigshafen, Germany. We are thankful to our co-workers Ralf Thomann, Victor Hugo Pacheco Torres, Marina Hagios, and Julia Hettenbach for their enthusiastic support. Donations of various glycidyl ethers from CVC Emerald Materials are gratefully acknowledged.

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ABBREVIATIONS PUR, polyurethane; PHU, polyhydroxyurethane; TPE, thermoplastic elastomer; TPHE, thermoplastic polyhydroxyurethane elastomer; FDM, fused deposition modeling; MH, mechanical hysteresis; SI, Supporting Information

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REFERENCES

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