Article pubs.acs.org/Macromolecules
Bottlebrush-Guided Polymer Crystallization Resulting in Supersoft and Reversibly Moldable Physical Networks William F. M. Daniel,† Guojun Xie,‡ Mohammad Vatankhah Varnoosfaderani,† Joanna Burdyńska,‡ Qiaoxi Li,† Dmytro Nykypanchuk,§ Oleg Gang,§,∥ Krzysztof Matyjaszewski,*,‡ and Sergei S. Sheiko*,† †
Department of Chemistry, University of North Carolina at Chapel Hill, Chapel Hill, North Carolina 27599-3290, United States Department of Chemistry, Carnegie Mellon University, Pittsburgh, Pennsylvania 15213, United States § Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, New York 11973, United States ∥ Department of Chemical Engineering, Columbia University, New York, New York 10027, United States ‡
S Supporting Information *
ABSTRACT: The goal of this study is to use ABA triblock copolymers with central bottlebrush B segments and crystalline linear chain A segments to demonstrate the effect of side chains on the formation and mechanical properties of physical networks cross-linked by crystallites. For this purpose, a series of bottlebrush copolymers was synthesized consisting of central amorphous bottlebrush polymer segments with a varying degree of polymerization (DP) of poly(n-butyl acrylate) (PnBA) side chains and linear tail blocks of crystallizable poly(octadecyl acrylate-stat-docosyl acrylate) (poly(ODA-stat-DA)). The materials were generated by sequential atom transfer radical polymerization (ATRP) steps starting with a series of bifunctional macroinitiators followed by the growth of two ODA-stat-DA linear-chain tails and eventually growing poly(nBA) side chains with increasing DPs. Crystallization of the poly(ODA-stat-DA) tails resulted in a series of reversible physical networks with bottlebrush strands bridging crystalline cross-links. They displayed very low moduli of elasticity of the order of 103−104 Pa. These distinct properties are due to the bottlebrush architecture, wherein densely grafted side chains play a dual role by facilitating disentanglement of the network strands and confining crystallization of the linear-chain tails. This combination leads to physical cross-linking of supersoft networks without percolation of the crystalline phase. The cross-link density was effectively controlled by the DP of the side chains with respect to the DP of the linear tails (nA). Shorter side chains allowed for crystallization of the linear tails of neighboring bottlebrushes, while steric repulsion between longer side chains hindered the phase separation and crystallization process and prevented network formation.
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INTRODUCTION Linear ABA block copolymers are a valuable class of polymeric materials capable of self-assembling into physical networks through microphase separation of the constituent polymer blocks.1−3 These networks are subject to intrinsic entanglements between their strands, which set a lower limit for the elastic modulus. By replacing the linear B segments with bottlebrushes, the strands disentangle, opening an opportunity for the creation of supersoft, reversible polymeric networks.4 Bottlebrush polymers are readily used as mesoscopic building blocks in the materials community and are comprised of three independently controlled structural parameters: the degree of polymerization (DP) of the backbone (nbb), the DP of the welldefined pendent side chains (nsc), and the number of backbone monomers per side chain, related to graft density (ng).5−10 Because of the high graft density of large pendent groups, bottlebrushes display a combination of extended but still flexible backbone conformations and large effective molecular diameters leading to disentanglement of their melts.4,11,12 By introducing cross-links into bottlebrush melts, one can create elastomers with exceedingly low moduli (down to the hundreds of Pa) while simultaneously allowing for high deformations and strong strain stiffening behavior.13 Moreover, steric repulsion © XXXX American Chemical Society
between the densely grafted side chains may strongly influence microphase separation behavior in ABA type copolymers and confine crystallization, thus giving additional control over both the network formation and its mechanical properties. This paper employs crystallization between linear tail segments to physically cross-link bottlebrush polymers to create supersoft, dry elastomers with the potential to be remolded and recycled through the melting of temporary crystalline junctions. In order to create these semicrystalline bottlebrush polymers, we turn to block copolymer bottlebrushes,14−24 specifically an ABA architecture (Figure 1) composed of two crystallizable linear chain A blocks and an amorphous bottlebrush segment (B middle block). Unlike the linear ABA counterparts,25−28 the bottlebrush block plays four distinct roles in controlling the network architecture formation. First, an increase in nsc effectively reduces the cross-link density of the disentangled network strands. Second, the side chains enhance the bending rigidity of the network strands favoring bridges versus loops of the B blocks, and hence promoting network formation. Third, steric repulsion between the densely grafted side chains of the Received: January 5, 2017
A
DOI: 10.1021/acs.macromol.7b00030 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules
brushes with A block tails of semicrystalline poly(octadecyl acrylate-stat-docosyl acrylate) and central B blocks of amorphous poly(n-butyl acrylate) bottlebrushes (poly[(ODAstat-DA)-b-(BiBEM-g-nBA)-b-(ODA-stat-DA)]) were synthesized using ATRP31,32 and then characterized for molecular and bulk network formation properties using DSC (differential scanning calorimetry), LB-AFM (Langmuir−Blodgett deposition atomic force microscopy), and shear rheology as a function of ODA-stat-DA content and bottlebrush molecular parameters. The specific chemical compositions of these materials are not essential to the phase separation behavior. However, ODA-statDA was chosen for the A blocks in order to control its melting temperature.
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Figure 1. Representation of soft network formation in ABA bottlebrush polymers. (a) Short side chains result in unconstrained crystallization, leading to ridged crystal scaffolds. (b) Intermediate side chain DPs result in association of neighboring A blocks, the formation of confined crystalline junctions, and soft bottlebrush networks. (c) Large side chains prevent association of A blocks, resulting in bottlebrush melts with isolated crystals.
EXPERIMENTAL SECTION
Materials. Octadecyl acrylate (ODA, 97%, Aldrich) was purified by dissolution in hexane and extraction four times with 5% aqueous NaOH. After drying the organic phase over anhydrous magnesium sulfate, the solution was passed through neutral alumina and solvent was removed under reduced pressure. (2-Trimetylsiloxy)ethyl methacrylate (HEMA-TMS, Scientific Polymer Products) and nbutyl acrylate (nBA, 99%, Acros) were purified by passing the monomer through a column filled with basic alumina to remove the inhibitor. Docosyl acrylate (DA), ethylene bis(2-bromoisobutyrate) (2f-BiB), and N′,N″-dioctadecyl-N′,N″-bis[2-(4-methoxy-3,5dimethyl)pyridylmethyl]ethane-1,2-diamine (DOD-BPED*) was synthesized according to procedures reported in the literature.33−35 Copper(I) bromide (CuBr, 99.999%, Aldrich), copper(I) chloride (CuCl, ≥99.995%, Aldrich), copper(II) bromide (CuBr2, 98%, Acros), copper(II) chloride (CuCl 2 , ≥99.995%, Aldrich), tris[2(dimethylamino)ethyl]amine (Me6TREN), ethyl α-bromoisobutyrate (EBiB), potassium fluoride (KF, 99%, Aldrich), tetrabutylammonium fluoride (TBAF, 1.0 M in THF, Aldrich), α-bromoisobutyryl bromide (98%, Aldrich), 2,5-di-tert-butylphenol (DTBP, 99%), triethylamine (TEA, ≥99%, Aldrich), and solvents were used as received without further purification. Methods. Synthesis of Difunctional Poly(HEMA-TMS). The general procedure is described below for the synthesis of P(HEMATMS)220. A 25 mL Schlenk flask was charged with ethylene bis(2bromoisobutyrate) (2f-BiB, 66.2 mg, 0.184 mmol), HEMA-TMS (20.0 mL, 91.9 mmol), dNbpy (0.1503 g, 0.3675 mmol), CuBr2 (8.2 mg, 0.037 mmol), and anisole (5 mL). The solution was degassed by three freeze−pump−thaw cycles. During the final cycle, the flask was filled with nitrogen, and CuBr (21.0 mg, 0.147 mmol) was quickly added to the frozen reaction mixture. The flask was sealed, evacuated, and backfilled with nitrogen five times and then immersed in an oil bath at 40 °C. Polymerization was stopped when conversion reached 43.9%, as determined by 1H NMR spectroscopy. From the molar ratio of the reacted monomer to initiator, the degree of polymerization (DP) was calculated to be 220. The reaction mixture was diluted with methylene chloride and passed through a neutral alumina column to remove the catalyst, and the solvent and the residual monomer were removed at high vacuum. The apparent molecular weight of the obtained poly(HEMA-TMS)220 determined by THF SEC using linear PMMA standards was Mn,app = 32,200 and Đ = 1.10. Poly(HEMA-TMS)444 was prepared using a similar procedure with the stoichiometric ratios outlined in Table 1. Synthesis of Macroinitiator. The general procedure is described below for the synthesis of poly[(DA50-stat-ODA50)-b-BiBEM220-b-
brush repel neighboring bottlebrushes, which in turn hinders crystallization of the linear A blocks, leading to reduced crosslink functionality and increased effective network strand length and preventing crystal percolation. Fourth, nsc is a primary control parameter for tuning the mass fraction of the crystalline phase (φw,A) for a given DP of the A block, as seen in eq 1 mA nA mA nA 1 φw,A ≈ ≈ ∼ mBA nbbnsc + mA nA mBA nbbnsc nsc (1) where mBA and mA are the monomer molar masses of nBA side chains and the ODA-stat-DA A blocks and nA is the DP of the A block. This work will build on past experience29,30 and focus on the reversible network formation properties afforded by the combination of ABA structures and bottlebrush architectures. Depending on the DP of side chains, we anticipate three distinct scenarios for network formation: (i) short side chains (high fraction of the A blocks) will allow for unconstrained crystallization of the A blocks, creating a hard crystalline scaffold effectively screening the mechanical properties of the bottlebrush B phase (Figure 1a), (ii) intermediate side chain sizes will allow for strong phase separation into discrete A domains, leading to soft percolated networks with reversible crystalline cross-links embedded in a continuous amorphous phase (Figure 1b), and (iii) steric repulsion between long side chains will hinder crystallization of the A blocks, leading to a decrease of cross-link functionality and eventually to un-crosslinked melts of bottlebrushes (Figure 1c). Because the first scenario is well noted and considered trivial, we aim at the effect of the side chains on the network formation of ABA bottlebrushes. Herein, we present the synthesis, characterization, and preliminary thermomechanical exploration of semicrystalline bottlebrush copolymers. A series of bottle-
Table 1. Experimental Conditions and Molecular Parameters for the Synthesis of P(HEMA-TMS)
a
sample
[M]:[2f-BiB]:[CuBr]:[CuBr2]:[dNbpy]
solvent
conv.a
Mn,appb
Mw/Mnb
poly(HEMA-TMS)220 poly(HEMA-TMS)444
500:1:2:0.8:0.2:2 1000:1:0.8:0.2:2
20% 20%
43.9% 44.4%
32,200 62,800
1.10 1.09
Conversion of HEMA-TMS measured by 1H NMR. bDetermined by SEC using linear PMMA standards. B
DOI: 10.1021/acs.macromol.7b00030 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules Table 2. Experimental Conditions and Molecular Parameters for the Synthesis of ABA Triblock Macroinitiator sample MI1 MI2 MI3 a
composition
[ODA]:[DA]:[poly(HEMA-TMS)]:[CuCl]:[CuCl2]:[DOD-BPED*] solvent
poly[(DA50-stat-ODA50)-b-BiBEM220b-(DA50-stat-ODA50)] poly[(DA105-stat-ODA105)-b-BiBEM220b-(DA105-stat-ODA105)] poly[(DA122-stat-ODA122)-b-BiBEM444b-(DA122-stat-ODA122)]
conv.a
Mn,appb
Mw/Mnb
411:411:1:4.9:0.5:5.4
55%
24.3%
49,600
1.47
800:800:1:6.7:1.3:8
59%
26.2%
54,400
1.50
800:800:1:7:1:8
51%
30.4%
63,400
1.64
Conversion of acrylate measured by 1H NMR. bDetermined by SEC using linear PMMA standards.
Table 3. Formulations Used for the Synthesis of the ABA Triblock Brush
a
sample
composition
[nBA]:[BiBEM]: [CuBr]:[CuBr2]:[dNbpy]
solvent
conv.a
SC-23 SC-34 SC-44 SC-110 SC-67-nAx2 SC-60-nAx2-nbbx2
poly[(DA50-stat-ODA50)-b-(BiBEM-g-nBA23)220-b-(DA50-stat-ODA50)] poly[(DA50-stat-ODA50)-b-(BiBEM-g-nBA34)220-b-(DA50-stat-ODA50)] poly[(DA50-stat-ODA50)-b-(BiBEM-g-nBA44)220-b-(DA50-stat-ODA50)] poly[(DA50-stat-ODA50)-b-(BiBEM-g-nBA110)220-b-(DA50-stat-ODA50)] poly[(DA105-stat-ODA105)-b-(BiBEM-g-nBA67)220-b-(DA105-stat-ODA105)] poly[(DA122-stat-ODA122)-b-(BiBEM-g-nBA60)444-b-(DA122-stat-ODA122)]
300:1:0.35:0.025:0.75 800:1:0.974:0.026:2 340:1:0.333:0.017:0.7 800:1:0.974:0.026:2 600:1:0.637:0.025:1.324 610:1:0.737:0.025:1.524
10% 10% 10% 10% 10% 10%
7.6% 4.3% 12.9% 13.8% 11.2% 9.9%
Conversion of acrylate measured by 1H NMR.
(DA50-stat-ODA50)]. A 25 mL Schlenk flask was charged with difunctional P(HEMA-TMS)220 (1.77 g, 0.0398 mmol), octadecyl acrylate (5.31 g, 16.4 mmol), docosyl acrylate (6.23 g, 16.4 mmol), CuCl2 (2.5 mg, 0.018 mmol), DOD-BPED* (0.170 g, 0.196 mmol), and o-xylene (19 mL). The solution was degassed by three freeze− pump−thaw cycles. During the final cycle, the flask was filled with nitrogen, and CuCl (17.7 mg, 0.178 mmol) was quickly added to the frozen reaction mixture. The flask was sealed, evacuated, and backfilled with nitrogen five times and then immersed in an oil bath at 90 °C. Polymerization was stopped when monomer conversion reached 24.3%, as determined by 1H NMR spectroscopy. From the molar ratio of the reacted monomer to initiator, the total degree of polymerization (DP) of each side block was calculated to be 100. Assuming equal incorporation (by mole) of ODA and DA in the side block, the composition was calculated to be poly[(DA50-stat-ODA50)-b-(HEMATMS)220-b-(DA50-stat-ODA50)]. The reaction mixture was diluted with THF, passed through a neutral alumina column to remove the catalyst, concentrated, and the product precipitated by addition to acetone and dried under a vacuum. A 100 mL round-bottom flask was charged with the product of the previous reaction (2.40 g, containing 3.75 mmol of HEMA-TMS repeat unit), KF (0.266 g, 4.50 mmol), DTBP (0.077 g, 0.38 mmol), and then dry THF (50 mL) was added under nitrogen. The reaction mixture was cooled down in an ice bath, followed by the injection of tetrabutylammonium fluoride (0.038 mL, 1.0 M in THF, 0.038 mmol) and subsequent dropwise addition of α-bromoisobutyryl bromide (1.0 g, 0.56 mL. 4.5 mmol) over the course of 20 min. The reaction mixture was then allowed to reach room temperature and was stirred for another 24 h. Afterward, the solids were filtered off and the mixture was precipitated by addition of the filtered solution to methanol/water (70/30) and the solids were redissolved in THF and passed through the column filled with basic alumina. The product was reprecipitated three times in methanol and dried overnight under a vacuum. The apparent molecular weight was determined by using linear PMMA standards in THF SEC: Mn,app = 49,600 and Mw/Mn = 1.47. The precursors (poly[(DA50-stat-ODA50)-b-(HEMA-TMS)220-b(DA50-stat-ODA50)]) of another two triblock ABA macroinitiator were prepared according to the above procedure using the stoichiometric ratios shown in Table 2, followed by modification with α-bromoisobutyryl bromide (using a stoichiometric ratio for the preparation of poly[(DA50-stat-ODA50)-b-BiBEM220-b-(DA50-statODA50)]). Synthesis of Copolymer Bottlebrush. The general procedure is described below for the synthesis of poly[(DA50-stat-ODA50)-b(BiBEM-g-nBA)220-b-(DA50-stat-ODA50)]. A 25 mL round-bottom flask was charged with MI1 (0.195 g, containing 0.326 mmol of initiating sites), nBA (14.0 mL, 97.8 mmol), CuBr2 (1.8 mg, 0.0081
mmol), dNbpy (0.100 g, 0.245 mmol), and anisole (1.6 mL). The solution was degassed by three freeze−pump−thaw cycles. During the final cycle, the flask was filled with nitrogen, and CuBr (16.3 mg, 0.114 mmol) was quickly added to the frozen reaction mixture. The flask was sealed, evacuated, and backfilled with nitrogen five times and then immersed in an oil bath at 70 °C. Polymerization was stopped when monomer conversion reached 7.6%, as determined by 1H NMR spectroscopy. From the molar ratio of the reacted monomer to initiator, the degree of polymerization (DP) of the side chain was calculated to be 23, assuming the initiation efficiency is 100%. The reaction mixture was diluted with THF, passed through a neutral alumina column to remove the catalyst, concentrated, and the product precipitated by addition to methanol and dried under a vacuum. The apparent molecular weight of the obtained triblock copolymer was determined by THF SEC using linear PMMA standards: Mn,app = 145,000 and Đ = 1.53. Five other triblock ABA brushes were prepared according to the above procedure using the stoichiometric ratios shown in Table 3. Synthesis of Linear Random Copolymer of Poly[(ODA50-statDA50). The general procedure is described below for the synthesis of poly[(DA50-stat-ODA50). A 10 mL Schlenk flask was charged with ethyl α-bromoisobutyrate (EBiB) (9.67 mg, 0.05 mmol), octadecyl acrylate (ODA) (1.00 g, 3.10 mmol), docosyl acrylate (DA) (1.17 g, 3.10 mmol), Me6TREN (11.5 mg, 0.05 mmol), CuCl2 (1.0 mg, 0.007 mmol), and anisole (5 mL). The solution was degassed by three freeze−pump−thaw cycles. During the final cycle, the flask was filled with nitrogen, and CuCl (3.7 mg, 0.035 mmol) was quickly added to the frozen reaction mixture. The flask was sealed, evacuated, and backfilled with nitrogen five times and then immersed in an oil bath at 45 °C. Polymerization was stopped when conversion reached 80.0%, as determined by 1H NMR spectroscopy. From the molar ratio of the reacted monomer to initiator, the degree of polymerization (DP) was calculated to be 100. The reaction mixture was diluted with methylene chloride, passed through a neutral alumina column to remove the catalyst, concentrated, and the product precipitated by addition to acetone and dried under a vacuum. Cross-Linking Procedure. A sample of SC-23 was molded into a 0.7 mm height, 8 mm diameter disk. The disk was exposed to 80 °C in a nitrogen atmosphere for a period of 24 h to induce homolytic cleavage of side chain bromine−carbon bonds, and subsequent random radical cross-linking of the samples into a permanent elastomer. Size Exclusion Chromatography (SEC) and Proton Nuclear Magnetic Resonance (1H NMR). The apparent number-average molecular weights (Mn,app) and dispersity (Mw/Mn) were measured by size exclusion chromatography (SEC). The SEC was conducted with a Waters 515 pump and Waters 2414 differential refractometer using PSS columns (SDV 105, 103, 500 Å) with THF as eluent at 35 C
DOI: 10.1021/acs.macromol.7b00030 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules Scheme 1. Synthesis of Triblock Molecular Brushes
Table 4. Structural Parameters and LB-AFM Size Analysis of Triblock Bottlebrush Polymers nsc,a NMR
nbb,b NMR
23 34 44 110
222 222 222 222
SC-67-nAx2
67
222
SC-60-nAx2-nbbx2
60
444
sample SC-23 SC-34 SC-44 SC-110
nA,c NMR
nBe
Lf (nm)
Wg (nm)
φw,A,h DSC
224 ± 3 5300 53.0 ± 0.4 23.9 ± 0.2 0.09 224 ± 3 7800 54.8 ± 0.5 37.0 ± 0.2 0.07 224 ± 3 10000 57.0 ± 0.5 55.9 ± 0.5 0.05 224 ± 3 25000 59.2 ± 0.6 89.1 ± 0.6 0.02 Increase of the Crystalline Block Length (nA) 2 × 210 241 ± 17 15000 64.5 ± 0.7 71.0 ± 0.3 0.08 Increase of Both Brush (nbb) and Linear (Crystalline) (nA) Segments 2 × 243 408 ± 4 27000 102.4 ± 0.9 68.9 ± 0.3 0.04 2 2 2 2
× × × ×
nbb,d AFM
100 100 100 100
Mn,i SEC (106)
Mn,j AFM (106)
Mn,k target (106)
0.15 0.24 0.36 0.50
0.7 2.3 1.8 2.7
0.8 1.0 1.3 3.3
0.41
3.0
2.1
0.39
3.4
3.6
a
DP of the side chains as determined by NMR measurement of monomer conversion. bDP of the bottlebrush backbone as determined by NMR measurement of monomer conversion. cDP of the semicrystalline tails as determined by NMR measurement of monomer conversion. dThe DP of the bottlebrush backbone determined as L/l, where l is the monomer length. eDP of the amorphous center B block. fLength of the amorphous B block as measured from AFM micrographs. gWidth of the bottlebrush B block as determined from AFM micrographs. hWeight fraction of ODA-statDA taken from targeted DPs of nsc and nA and measured values of nbb (see Table 1) calculated as eq 1. iNumber-average molecular weight as determined by SEC. jNumber-average molecular weight as determined by the LB-AFM technique. kNumber-average molecular weight as calculated from DPs of the bottlebrush samples. °C and at a flow rate of 1 mL min−1. Linear PMMA standards were used for calibration. Polymerization was monitored by 1H NMR spectroscopy using a Bruker Advance 300 MHz NMR spectroscope with CDCl3 as a solvent. Monomer conversion was calculated from the decrease of the monomer peak area relative to the peak areas of the internal standards. Langmuir−Blodgett Deposition Atomic-Force Microscopy (LBAFM). Known volumes (30−100 μL) of bottlebrush polymer in chloroform solutions (0.03−0.06 mg/mL) were deposited dropwise onto the water−air interface of a KSV 3000 series Langmuir−Blodgett trough. The samples were then compressed at a rate of 20 mm/min until monolayer formation was observed. LB films were transferred onto cleaved mica substrates at a constant surface pressure of 0.5 mN/ m and a controlled transfer ratio. AFM micrographs were taken in PeakForce QNM mode using a multimode AFM (Brüker) with a NanoScope V controller. Silicon probes were used with a resonance frequency of 50−90 Hz and a spring constant of ∼0.4 N/m. In-house developed computer software was used to analyze the AFM images for length and width distributions. Collections of over 1000 molecules were analyzed to ensure standard deviations of the mean remained below 10%. Molecular weights of samples were then calculated by using the known mass of polymer deposited onto the trough surface and the known trough surface area during film transfer to find the mass per unit area. This was then combined with the known number of
bottlebrush molecules per unit area taken from AFM micrographs to find the number-average molecular weight. Thick films for AFM were produced from 10 mg/mL in chloroform solutions deposited in 0.2 mL aliquots onto cleaved mica surfaces. The films were allowed to dry in a closed Petri dish at room temperature before being imaged. Differential Scanning Calorimetry (DSC). The transition temperatures and heats of melting and crystallization were measured using a TA Q200 series DSC. Measurements were performed in a range from −100 to 100 °C using heating and cooling rates of 3 °C/min with samples sizes between 4 and 6 mg. The results are summarized in Table 5. Small/Wide-Angle X-ray Scattering. Measurements were conducted at Brookhaven National Lab, Center for Functional Nanomaterials, using a Bruker Nanostar U instrument equipped with Vantec 2000 area detectors and a copper rotating anode as a source. Scattering patterns from the samples were measured at room temperature, and the collection time was set as 30 s for all samples. Shear Rheology. Mechanical properties of the ABA bottlebrush triblock copolymers were measured using an ARES-G2 shear rheometer from TA Instruments with 8 mm steel parallel plate geometries. For the stress relaxation curves, samples were loaded and their thermal histories were erased by allowing all samples to sit at 70 °C for a period of 30 min before being compressed and trimmed into sample disks of 0.7 mm height by 8 mm diameter. The samples were D
DOI: 10.1021/acs.macromol.7b00030 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules Table 5. Thermal Parameters of the Bottlebrush Samples from DSC and Apparent Shear Moduli sample
φw,A,a target
φw,A,b DSC
Tm,c DSC (°C)
Tc,d DSC (°C)
Tg,e DSC (°C)
ΔHmf (J/g)
ΔHcg (J/g)
G0h (Pa)
SC-23 SC-34 SC-44 SC-110 SC-60-nAx2-nbbx2 SC-67-nAx2
0.09 0.07 0.05 0.02 0.04 0.08
0.09 0.08 0.05 0.02 0.04 0.07
55 55 53 53 53 53
41 41 41 42 41 41
−47 −47 −47 −49 −47 −49
7.8 6.6 4.2 1.8 3.4 5.5
8.5 7.3 5.3 2.1 4.0 6.1
13000 5000 1900 0 NA 2900
a Weight fraction of ODA-s-DA taken from DPs of nsc and nA and measured values of nbb (see Table 1) calculated as eq 1. bWeight fraction of ODAs-DA taken from the ratio of ΔHm/ΔH0m, where ΔH0m = 84 J/g is the heat of melting of in house synthesized free ODA-stat-DA tails. cTemperatures of melting from DSC curves. dTemperatures of crystallization from DSC curves. eGlass transition temperature taken from the middle of the heating transition from the DSC curves. fHeats of melting calculated from the area under the melting peak of the DSC curves. gHeats of crystallization calculated from the area under the crystallization peak of the DSC curve. hThe apparent shear modulus of the crystallized sample taken from stress relaxations after 3 h of relaxation at 25 °C.
then rapidly cooled to the experimental temperature, allowed to equilibrate at temperatures of 70 or 25 °C for 30 min, subjected to a 5% strain, and measured over a period of 3−24 h to ensure stable stress values or complete relaxation to 0 Pa of stress. The thermal ramps were executed using an identical loading procedure to the one described above followed by equilibration at 70 °C. The samples were then subjected to 10% strain at a frequency of 1 rad/s while cooling down to 25 °C at a rate of 1 °C/min.
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RESULTS AND DISCUSSION The general synthetic procedure for the preparation of bottlebrush copolymers is shown in Scheme 1. Poly[(ODAstat-DA)-b-(BiBEM-g-nBA)-b-(ODA-stat-DA)] brushes were synthesized by grafting nBA from poly[(ODA-stat-DA)-bBiBEM-b-(ODA-stat-DA)] macroinitiator. The triblock backbone macroinitiator was synthesized via ATRP of HEMA-TMS using a difunctional initiator, ethylene bis(2-bromoisobutyrate) (2f-BiB), followed by chain extension of ODA and DA from the resultant difunctional poly(HEMA-TMS) and esterification of the macroinitiator precursor with α-bromoisobutyryl bromide. As outlined in the Experimental Section, the structural parameters (nsc, ng, nbb) of the bottlebrush B block (collectively called nB = nbb(1 + nsc/ng)) and DP of the crystalline A blocks (nA) were precisely tuned by changing the initial monomer/ initiator molar ratio and the final conversion of monomers in polymerization. 1H NMR analysis combined with SEC was used to identity the targeted values of nsc, nbb, nA, and molecular weight (Mw). The results of molecular characterization are summarized in Table 4. From SC-23 to SC-110, when the DP of the side chain increases from 23 to 110, the weight fraction of crystallizable tails (ODA-stat-DA) correspondingly decreases from 9 to 2% (Table 5). The length of the A block for SC-67nAx2 is approximately 2 times larger than that of SC-34. Therefore, to have a similar weight fraction of crystallizable side blocks, the DP of side chains for SC-67-nAx2 is almost twice as large as that of SC-34. Also, although having a similar length of side chain and side blocks as SC-67-nAx2, SC-60-nAx2-nbbx2 has a significantly lower weight fraction of crystallizable blocks due to a much longer central block. In addition, molecular imaging by AFM was employed to find the length (L), 2D width (W) distributions, and true molecular weights of the samples (Figure 2 and Table 4). These experiments display excellent agreement between the measured lengths (L) (calculated nbb) and the targeted nbb, as well as good agreement between the targeted values of nsc and the widths (W) which are expected to scale as nsc = nsc,w/ĐX ∼ W/ ĐX, where nsc,w is the weight-average DP of the side chains and ĐX is the dispersity of the side chains (Supporting Figure S1).36
Figure 2. AFM of 2D crystallized bottlebrush copolymers. (a) AFM height micrographs of single ABA bottlebrush macromolecules linked by crystalline junctions of the poly(ODA-stat-DA) linear tails (A block). (b) AFM height micrographs of LB monolayers of ABA bottlebrush copolymers with varying PnBA side chain lengths. Inset images represent close-ups of crystalline junctions. (c) AFM height micrographs of thick films of the corresponding ABA copolymers at room temperature. The crystalline domains appear as bright elevated patches. (d) The number of crystalline junctions connecting at least two bottlebrushes decreases with nsc. Black squares (□) represent samples with nbb and nA of 224 and 2 × 100, respectively, and the red triangle (△) represents a sample with nbb and nA of 408 and 2 × 243, respectively.
Bottlebrush molecules possess large molecular dimensions which can be readily resolved by AFM to find the accurate number of molecules per unit area in dense LB monolayers (Figure 2b). Combining this knowledge with the known mass of polymer deposited on the known area of the LB trough allows for the accurate determination of the number-average molecular weight (Mn,AFM). These MW values obtained from E
DOI: 10.1021/acs.macromol.7b00030 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules
Figure 3. DSC of bottlebrush copolymer samples. (a) A representative DSC curve for sample SC-60-nAx2-nbbx2 displaying freezing peak, melting peak, and glass transition. (b) Relative heats of melting as ΔHm/ΔH0m, where ΔH0m is the heat of melting for ODA-s-DA (84 J/g). The values display a nearly linear relation with the targeted weight fraction of ODA-s-DA. Black squares (□) represent samples with nbb and nA of 224 and 2 × 100, respectively, and differing side chain lengths (SC-23, 24, 44, 110). The red triangle (△) represents a sample with nbb and nA of 408 and 2 × 243, respectively (SC-60-nAx2-nbbx2). The blue circle (○) represents a sample with nbb and nA of 241 and 2 × 210, respectively (SC-67-nAx2).
Figure 4. Rheological characterization of select bottlebrush copolymers. (a) Shear relaxation modulus, G0(t), as a function of time at 25 °C (below Tc) of samples SC-23 (black), SC-34 (purple), SC-67-nAx2 (red), SC-44 (blue), and SC-110 (green) displaying fast relaxation (inset) followed by plateaus at 13000, 5000, 2900, 1900, and 0 Pa, respectively. This indicates the presence of physically cross-linked networks. Inset: log−log plot of G0(t) at short times showing a power-law relaxation G0(t) ∼ t0.42±0.06 consistent with Rouse-like behavior. (b) The relaxation modulus, G0(t), of the same samples as a function of time at 70 °C (above Tm). The curves display rapid relaxation to zero value, indicating lack of network formation. (c) Oscillatory storage and loss moduli (G′ and G″) of sample SC-23 (1 rad/s, 5% strain) displaying sharp increases at a crystallization temperature of T = 41 °C. (d) The data points (□) corresponding to the normalized apparent relaxation modulus G0/nA2 taken from part a at t = 3 h decrease with the DP of the side chains (nsc). Note that SC-110 does not form a network (G = 0). The blue dashed line represents the expected trend in G0/nA2 when nx = nbb taken from eq 2 substituting eqs 3 and 4. As can be seen, the networks display a stronger dependence on nsc than expected. The red dashed line represents the fit of the experimental G0/nA2 taken from eq 2 substituting eqs 4 and 6 assuming that the cross-link density decreases with increasing nsc.
as seen in Figure 2b. Sample SC-23 displays a seemingly percolated network of crystalline junctions, while sample SC110 shows isolated crystals with few junctions. To reduce the effect of the substrate and approach the bulk morphology, thick films of samples SC-23, SC-44, and SC-110 were prepared by drop casting from chloroform solutions (Figure 2c). Sample SC-23 with the shortest side chains displays large and numerous crystalline domains. Despite the apparently high fraction of the crystallites, they do not seem to merge or percolate, as sketched in Figure 1b. SC-44 displays a similar density of crystallites but with smaller domains, and, as such, larger inter-cross-link distances suggesting a softer network. In
the LB-AFM analysis compare favorably with the targeted molecular weights (Table 4). Besides the MWD analysis, molecular imaging by AFM was also used to qualitatively verify network formation due to crystallization of the A tail blocks. For this purpose, we have prepared samples with different surface coverages including single molecules (Figure 2a), dense monolayers (Figure 2b), and thick films (Figure 2c). Samples of single molecules prepared by spin-casting from a dilute solution clearly display end-chain association of bottlebrush blocks (Figure 2a). The dense monolayers display large numbers of crystalline junctions which decrease with increasing nsc and decreasing φw,A (eq 1), F
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displaying slope near −0.5 and below). The apparent shear modulus of the network (G0) decreases with increasing nsc and decreasing φw,A displaying a relatively high modulus in the 104 Pa range for SC-23 and dropping down to zero network modulus for SC-110. It should be noted that sample SC-110 still displays obvious crystallization transitions in DSC measurements but no network formation which is likely due to side chain steric hindrance preventing contact between the crystalline tails of neighboring molecules. When taken above the melting temperature, all samples decay to a modulus of zero within a relatively short time span, displaying the lack of any permanent cross-linking mechanisms within the samples. Figure 4c displays a freezing transition for sample SC-23. There is a sharp increase in both the loss and storage modulus between 44 and 41 °C indicative of a crystallization transition and showing a relatively small change resulting in a cross-linked modulus on the order of 0.02 MPa (more thermal ramps are presented in Supporting Figure S3). From the data in Figure 4a,b, it can be seen that this polymer architecture is capable of creating stable but reversible physical networks with moduli below 0.1 MPa. The equilibrium apparent moduli normalized by the DP of the crystalline segments (G0/nA2) were then plotted against their respective nsc values (□, Figure 4d). Note that the modulus of bottlebrush networks (even when measured at small strains ε → 0) is affected by strain-stiffening due to the extended conformation of backbones in the bottlebrush strands. Due to finite extensibility of the backbone, this so-called apparent modulus is higher than the structural modulus and described by the Dobrynin equation37
contrast, the SC-110 film (on the right) displays only small areas of crystalline A phase and likely represents the third scenario in Figure 1c, where long side chains prevent association of A blocks, resulting in low cross-linking instances and failure of network formation. The 2-D images do not represent the true bulk organization of ABA networks and provide only a qualitative proof of network formation. Nevertheless, we plotted the areal density of the cross-link junctions versus the DP of the side chains (Figure 2d), which shows a progressive decrease of the number of junctions per unit area of the ABA monolayers. This observation is attributed to the enhancement of steric repulsion with the side chain length, which hinders the association of the A blocks. From AFM, it can be seen that the DP of the side chains plays a strong role in the formation of intermolecular crystallization and thus the potential cross-linking of bulk networks. While AFM gives a qualitative measure of interbrush crystallization, it is also necessary to quantify the total crystallinity of the samples, i.e., both intermolecular (crosslinking) and intramolecular (non-cross-linking). To quantify the true crystallinity of the samples, the degree of crystallization was measured by DSC as a function φw,A, which is directly related to the side chain DP as φw,A ∼ 1/nsc (eq 1). Figure 3a displays a representative DSC curve of sample SC-60-nAx2nbbx2. Despite the relatively low amount of crystalline phase, the sample displays a very clear crystallization and melting transitions. The values of the melting temperature Tm ranged from 53 to 55 °C, which matches well with the measured value for free ODA-stat-DA chains with DP = 103 (53 °C). The DSC analysis is summarized in Table 5. Figure 3b displays the relative heats of melting versus the calculated weight fraction of ODA-stat-DA and shows a linear relation. The close match between DSC and estimations of φw,A predicted on the basis of the targeted structures provides strong evidence for the accurate control of the targeted molecular parameters. This also gives evidence that the nsc plays a strong role in regulating network formation without preventing intramolecular crystallization. In other words, the close match between the calculated values φw,A and those values measured from DSC show that all samples fully crystallize, but only those with shorter side chains can form intermolecular junctions and networks. The temperatures of melting (Tm) clearly group between samples; however, there is an increase in the Tm for the SC-23 and SC-34. It may indicate formation of larger crystals due to the weaker effect of the shorter side chains on the crystallization of the A tails. The increase of crystallinity with decreasing length of the side chains was corroborated by X-ray (Supporting Information Figure S2). With the successful verification of molecular association and crystallization of the ABA bottlebrush copolymers, the network formation properties were assessed using shear rheology. To prove the formation of networks, samples with different side chain DPs were subjected to stress relaxation experiments at both 25 °C (well below the crystallization temperature of 41 °C seen in Table 5) and above the melting temperature at 70 °C. All samples were allowed to crystallize and melt at the respective temperatures for a period of 30 min to ensure consistent network formation conditions. Parts a and b of Figure 4 show stress relaxations of the bottlebrush copolymers measured below and above the crystallization and melting temperatures, respectively. At 25 °C (Figure 4a), the polymers display a constant network modulus after a brief period of Rouse-like and transition regime relaxations (figure inset
G0 =
G (1 + 2(1 − β)−2 ) 3
(2)
where G is the structural modulus determined by the cross-link density and β = ⟨Rin2⟩/Rmax2 is the strand elongation ratio defined as the ratio of the mean square end-to-end distance of a network strand in the as prepared elastomer (⟨Rin2⟩) and the square of the strand contour length (Rmax2). For bottlebrush networks with nsc ≫ 1 and the mass fraction of the crystalline phase φw,A (eq 1), the structural modulus and the strand elongation ratio depend on the architectural parameters as13 G≅
mA nA ⎞ ρRT ρRT ⎛ (1 + φw,A ) ≅ ⎜1 + ⎟ mBA n x nsc mBA n x nsc ⎝ mBA nbbnsc ⎠ (3)
β≅
1/2 v1/2 nsc l 3/2 n x
(4)
where ρ is the mass density of the polymer, RT is the molar thermal energy, mBA is the molar mass of the amorphous monomer, mA is the averaged molar mass of the ODA-stat-DA monomers, l ≅ 2.5 Å is the length of the (meth)acrylate monomers through the main chains, and v ≅ 195 Å is the volume of the nBA monomer. In this model, the A blocks act simply as cross-linkers and each ABA molecule is assumed to span two multifunctional A block junctions, making nx = nbb. An estimation for this trend can be seen in Figure 4d as the blue dashed line representing G0/nA2 from eq 2 substituting eqs 3 and 4. It became obvious that the materials possessed a much stronger dependence for both G0 and β, which is ascribed to a decrease of the cross-link density due to an increase of the side chain length. As discussed in the Introduction (Figure 1), stronger steric repulsion between longer side chains hinders G
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original shape upon heating back to 55 °C. This behavior gives evidence that such systems will be suited to future applications as shape memory biomedical implants where both the crystalline and amorphous moduli match those of surrounding tissues.
association of A tails from neighboring ABA bottlebrushes. This in turn effectively leads to a decrease in the cross-link functionality and an increase in the average network strand length. Strong evidence of this phenomenon can be seen in Figure 2a,c and Figure 4a where sample SC-110 displays small disconnected crystallite in monolayer films, sparse crystalline junctions in bulk films, and no network modulus below the sample Tm. New equations assuming that the degree of polymerization of the network strand increases with nsc as nx ≈ nsc/nA were derived as G≅
β≅
ρRTnA ⎛ mA nA ⎞ n 2 1+ ⎟ ∼ A3 2⎜ mBA nbbnsc ⎠ mBA nsc ⎝ nsc
(5)
v1/2 nA l 3/2 nsc1/2
(6)
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CONCLUSIONS A series of ABA triblock copolymers with semicrystalline tails and central amorphous bottlebrush blocks was synthesized. The materials were characterized by a variety of techniques to confirm their structural parameters. All of the samples display crystallization of the poly(ODA-stat-DA) tails, resulting in the formation of physical networks. The bottlebrush architecture plays a dual role in determining the mechanical properties of these networks. First, steric repulsions between the densely grafted side chains facilitate disentanglement of the network strands, eliminating the typical lower limits for network modulus. Second, the same repulsions confine the crystallization process, preventing percolations of a crystalline scaffold. The absence of entanglements paired with controlled crystallization allows for the formation of remoldable physical networks with supersoft moduli, orders of magnitude below both typical semicrystalline networks and amorphous elastomers. The apparent shear modulus G0 of these networks decays with nsc, until the absence of network formation when the nsc becomes comparable to nA, effectively isolating the A blocks and preventing cross-linking. We discovered the system is best modeled by postulating that the nx is controlled by changing nsc rather than simply being determined by nbb, as would be the case in an ideal and fully cross-linked system. The ability of the bottlebrushes in these physical networks to prevent percolation of crystallites and control cross-linking density by tuning architectural parameters enables the creation of supersoft elastomers with the capacity for reversible remolding and shape memory. Further study into the optimization of such systems could lead to new generations of biocompatible thermosensitive materials, recyclable elastomers, and removable pressure sensitive adhesives.
These equations account for the change in nx as nsc (φw,A) changes as well as an increase in the expected values of β due to repulsion between phase-separated components of the bottlebrush copolymers causing extended backbone conformations. The final expression in eq 2 substituting eqs 5 and 6 accurately follows the experimental trend (red dashed line in Figure 4d). This corroborates that the physical cross-linking of these bottlebrush copolymers is influenced by the steric repulsion created by the side chains near the crystalline tails, which is also evidenced by the lack of plateau modulus in SC-110 despite its crystallization. These materials not only form supersoft physical networks but also show the potential for shape memory behavior. In general, shape memory results from counteraction between a covalent permanent network and a physical temporary network. While the permanent network controls the original shape, an external stress is applied to deform the sample into a programmed shape and temporary cross-linking elements such as crystallites are then introduced to fasten the programmed shape in place. Subsequent removal of the temporary cross-links, e.g., by melting crystallites, allows the permanent network to reassert itself and the permanent shape to restore. Usually, covalent networks are dense and, therefore, require significant crystallinity to fix the programmed shape. However, the soft bottlebrush networks require much less crystallinity to allow for a shape memory effect. Figure 5 illustrates a shape memory experiment for which polymer SC23 was thermally cross-linked at 80 °C for 24 h, resulting in a covalent chemical network with a shear modulus on the order of 200 Pa. The sample was bent to a 47° angle at 55 °C, then fixed by quenching to 22 °C, and subsequently returned to its
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ASSOCIATED CONTENT
* Supporting Information S
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b00030. Information on the characterization and rheology of bottlebrush materials (PDF)
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected]. *E-mail:
[email protected]. ORCID
Krzysztof Matyjaszewski: 0000-0003-1960-3402 Sergei S. Sheiko: 0000-0003-3672-1611 Notes
The authors declare no competing financial interest.
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Figure 5. Supersoft shape memory. A shape memory cycle using a rectangular rod of covalently cross-linked SC-23 polymer which is bent to a 47° angle at an elevated temperature of 55 °C (above ODA melting) and then subsequently fixed by quenching to 22 °C. Upon heating to 55 °C, the sample recovers to its original rod-like shape.
ACKNOWLEDGMENTS The authors gratefully acknowledge funding from the National Science Foundation (DMR 1407645, DMR 1501324, DMR 1436219, and DMR 1436201) and from Becton Dickinson H
DOI: 10.1021/acs.macromol.7b00030 Macromolecules XXXX, XXX, XXX−XXX
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