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Article Cite This: ACS Omega 2018, 3, 6945−6954
Influence of Phase Separation on Performance of Graft Acrylic Pressure-Sensitive Adhesives with Various Copolyester Side Chains Yanjiao Wang,† Feiyin Weng,† Jiaxu Li,† Lei Lai,† Wei Yu,*,‡ Steven John Severtson,*,§ and Wen-Jun Wang*,†
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†
State Key Laboratory of Chemical Engineering, College of Chemical and Biological Engineering, Zhejiang University, 38 Zheda Road, Hangzhou, Zhejiang 310027, China ‡ Advanced Rheology Institute, School of Chemistry and Chemical Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, China § Department of Bioproducts and Biosystems Engineering, University of Minnesota, 2004 Folwell Avenue, Saint Paul, Minnesota 55108, United States S Supporting Information *
ABSTRACT: Acrylic pressure-sensitive adhesives with various polyester side-chain lengths were synthesized to investigate the effect of branching on phase separation and polymer mechanical performance. The polyester macromonomers (MMs) were produced through ring-opening co-polymerizations of L-lactide (L-LA) and ε-caprolactone (ε-CL) initiated with 2hydroxyethyl methacrylate (HEMA), which provides the polyester chains with terminal vinyl groups. By varying the HEMA content, a range of MM chain lengths constructed from L10C4 (five L-LA and four ε-CL units) to L100C40 were obtained at a constant monomer mole ratio. Copolymerization of 2-ethylhexyl acrylate and acrylic acid with these MMs at constant mass composition provided a series of comb copolymers consisting of acrylic backbones with polyester branches of various chain lengths. Characterization of thin films cast from the polymers using thermal analysis and scanning probe microscopy showed a transition from a homogeneous phase to the formation of distinct microphases with increasing branching chain lengths. Rheological analysis of the linear viscoelastic responses was also used through small-amplitude oscillatory shear, and dynamic master curves were constructed by time− temperature superposition. The rheological data were also consistent with phase separation for the longer side-chain lengths of L50C20 and L100C40. The extra elastic contribution at low frequency and the temperature dependence of aT both show obviously effect of separated phases. Performance testing of polymer films showed that the chain extension resulted in a significant increase in both peel strength and shear resistance, which was accompanied by a modest decrease in film tackiness. The results demonstrate that tailoring branch chain structures provide a promising means for controlling the properties of the high-biomass content adhesive polymers.
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INTRODUCTION
structures. Although such copolymers tend not to possess adhesive characteristics,10,11 it was shown that the ratio between MM components could be used to control physical properties and important adhesive performance characteristics of the hybrid polymers. The role played by MM length in determining properties and the tendency of the lactide-based components to segregate are not evident. Such microstructural details are likely central to the successful application of these materials. The performance of a PSA is commonly gauged with three mechanical tests; tack force (tack), peel strength (peel), and shear resistance time (shear).12 As would be expected, changes
In recent years, there has been a great interest in the development of sustainable materials generated using renewable biomass. Efforts to replace monomers derived from fossil fuels are especially prevalent in the area of pressure-sensitive adhesive (PSA).1,2 Various renewable materials have been used as wholly or in-part substitutes for PSA components, including starch,3 polymerized soybean oil,4−6 and monomers, such as Llactide (L-LA), isosorbide, and menthyl acrylate.7−9 Amongst the more promising approaches is the generation of acrylic hybrid polymers. Previously, the synthesis, characterization, and performance of adhesive polymers possessing acrylic-based backbones and high-polyester-content short branches were reviewed.1,2 The short branches are introduced during synthesis as macromonomers (MMs) composed of acryliccapped L-LA and ε-caprolactone (ε-CL) linear chain © 2018 American Chemical Society
Received: April 17, 2018 Accepted: June 12, 2018 Published: June 26, 2018 6945
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ACS Omega that enhance cohesive strength of an adhesive film concurrently reduce its ability to flow and wet surfaces, i.e., shear resistance tends to be inversely related to both tack and peel performance. All performance properties are influenced by characteristics of the adhesive polymer, such as its molecular weight, monomer ratio, and monomer chain structure, as well as microstructure.13,14 Also, important in determining adhesive performance is the degree of cross-linking present, both chemical and physical.15−20 The influence of physical cross-linking is apparent in thermoplastic or hotmelt PSA. Traditional hotmelt is primarily composed of base polymer, tackifying resin, and processing oil (plasticizer). Base polymers are commonly block copolymers consisting of isoprene (I) or butadiene (B) in combination with styrene (S), such as the triblock copolymers SIS and SBS. The microphase separation of block components forms a repeating nanostructure consisting of hard glassy PS regions, which are often described as physical cross-links, and soft rubbery PI or PB regions. These microphases can be targeted with additives such as tackifying resin or diblock copolymer, allowing for the manipulation of the microstructure and adhesive properties.21−26 The adhesive polymers discussed in this work are branched copolymers. The use of pendant groups that are distinctly different (chemically or structurally) from the backbone provides another means for achieving novel property combinations.27−29 Previously, it was demonstrated that adhesive properties could be finely tuned through changes in low-molecular-weight pendant group compositions. Extending these chains produces long side chains and the possible aggregation of microstructures, as found for linear block copolymers. Given that the performance of PSA is dependent on the balance between elastic and viscous behavior, great insights are provided through their rheological characterization. Rheological properties of branched copolymers can be tailored by changing the side-chain length and composition. This is demonstrated primarily in the form of rheological data collected using small-amplitude oscillatory shear, an approach that has proven to be quite useful in detecting such effects due to its high sensitivity in the linear viscoelastic region.30 The viscoelastic properties of branched polymer melts are often described through various structural models, such as bottlebrushes,31 combs,32,33 stars,34,35 and so on. Bottlebrush polymers are those possessing a high density of branches, resulting in higher rigidities compared to those of linear polymers of similar molecular weights. The dynamic mechanical behavior of such polymers is governed by relaxation modes for both the arms and backbone,31,36 with the high-frequency behavior being dominated by the shorter relaxations of the arms and low-frequency behavior by the longer relaxations of the backbone. If the arm and backbone are both unentangled, their relaxation spectra are similar to the Rouse relaxation spectrum.31 For systems with entangled arms and backbone, two rubbery plateaus can be found in the plots of dynamic storage moduli (G′).37 For the rheological analysis of phase-separated systems, G′ values increase at low frequencies due to the extra elastic contribution of separated phases. The phase separation is often confirmed using the combination of rheological data, microscopy, and thermal analysis.38,39 It is well documented that a low elastic modulus is required for a polymer to demonstrate self-adhesive behavior, e.g., the Dahlquist criterion.40 This can be attained by increasing
entanglement molecular weight (Me) or lowering chain entanglement density. Introducing comb or bottlebrush structures in PSA is expected to increase Me. However, PSA with a low elastic modulus often possesses poor peel performance. Peel strength tends to correlate with the loss modulus (G″), which is indicative of enhanced energy dissipation along with strain hardening. Increasing G″ can be accomplished by broadening the glass transition temperature (Tg) range of a PSA or increasing its degree of cross-linking, which also enhances the creep resistance. The drawback of chemical cross-linking is that it strongly impacts the tack performance of PSA. Herein, the comb PSA having acrylic backbones and copolyester MM side chains is focused. A range of MMs with various chain lengths were used at a constant weight composition in the overall polymeric structures to examine the influence of branching on aggregational, rheological, and mechanical properties of the polymers. It was expected that the polyester MMs would have a poor compatibility with the acrylic backbone above their Mes,41 leading to physical crosslinking/ phase separation. Given that shorter block lengths are more favorable for phase separation,42 free-radical copolymerization of acrylates and copolyester MMs was used to produce the PSA with broad molecular weight distribution (Đ). It was expected that the low Me (for tack), broad Tg (for peel), and physical cross-linking (for peel and shear) of the PSA would provide good adhesive performance. As will be discussed, the results indicate that increasing the length of such pendent branch structures provides for a significant enhancement of adhesive polymer properties and the covalent linking of immiscible molecular chains to the backbone may provide greater stability and control over morphology and performance compared with blended and blocky materials.43,44
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EXPERIMENTAL SECTION Materials. Catalyst tin(II) 2-ethylhexanoate, Sn(Oct)2, (>95%), 2,2′-azobis(2-methylpropionitrile), AIBN (>98%), 2-hydroxyethyl methacrylate, HEMA, (>97%), ε-CL (>97%), acrylic monomers 2-ethylhexyl acrylate (2-EHA), acrylic acid (AA), methyl acrylate (MA), and solvents (ethyl acetate and tetrahydrofuran) were purchased from Sigma-Aldrich. L-LA (>99%) was generously provided by Shenzhen Bright China Industrial Co., Ltd. Deuterated chloroform (CDCl3) used to dissolve MMs and generated copolymers for 1H and 13C NMR analysis was purchased from J&K Chemicals. Synthesis of MMs and Branched Copolymers. Macromonomers of various chain lengths were synthesized from HEMA, L-LA, and ε-CL. The L-LA and ε-CL were combined at a constant molar ratio of 5:4, respectively, for various ratios of HEMA, mechanically stirred, and heated to 140 °C under a N2 atmosphere. Sn(Oct)2 (0.5 mol % based on monomer) was then introduced through a septum with a syringe to initiate polymerizations. Reactions were maintained at 140 °C for 4−6 h under constant mechanical stirring, with required reaction times increasing with lower HEMA molar ratios (i.e., larger MM chain lengths). Products were allowed to cool to ambient temperature prior to characterization and use. For the synthesis of the acrylic polymer, 2-EHA and AA were mixed with ethyl acetate to make a 50 wt % solution and then introduced into a round-bottomed flask equipped with a mechanical mixer. The solution was purged with N2 for 5 min and immersed in a 65 °C oil bath. Polymerization was initiated by charging AIBN dissolved in 2 mL of ethyl acetate. After 10 6946
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ACS Omega Scheme 1. Synthesis of Macromonomers (MM) and Acrylic Copolymers
Table 1. Characterization Results for Polyester MMs with Various Chain Lengthsa x (%)b
Mn (kg/mol)
MM
HEMA
L-LA
ε-CL
theor.
NMRc
GPCd
Đd
Tge (°C)
Fvinylf (%)
L10C4 L20C8 L50C20 L100C40
∼100 ∼100 ∼100 ∼100
∼100 ∼100 ∼100 ∼100
∼100 ∼100 ∼100 ∼100
1.30 2.48 6.01 11.89
1.42 2.59 6.23 12.47
2.04 2.13 5.17 7.68
1.65 1.70 1.76 1.66
−31.6 −24.8 −20.6 −15.7
∼100 ∼100 99 97
Ring-opening copolymerization reactions initiated with HEMA, catalyst = Sn(Oct)2, T = 140 °C, t = 4 h for L10C4 and L20C8, 5 h for L50C20, and 6 h for L100C40. bx = the conversion of monomers determined using 1H NMR. cDetermined using 1H NMR. dDetermined using GPC in tetrahydrofuran (THF), calibrated with PS standards. eDetermined by DSC. fFvinyl = percentage of the MM chain structures possessing a reactive terminal double bond.
a
temperatures Tg were gauged on a TA Instruments Q200 differential scanning calorimeter. Subsequent to thermal conditioning, samples were heated at a rate of 10 °C/min under a nitrogen environment over a temperature range of −75−150 °C to locate transitions. Surface morphologies of polymer samples were characterized using a Bruker Multimode 8 scanning probe microscope (SPM). Films of approximately 25 μm thickness were spincoated onto silicon wafers and conditioned for 24 h at 23 °C prior to testing. The SPM images were collected in tapping mode under ambient conditions. The linear viscosity rheological properties of all samples were gauged using a HAAKE RS6000 Rotational Rheometer equipped with 20 mm diameter parallel plates. Circular samples of approximately 1 mm thickness were obtained, filling a Teflon mold with the polymer dissolved in ethyl acetate and evaporating the solvent slowly. Frequency sweeps were carried out over a range of 0.01−40 Hz in oscillatory shear at a strain of 1% for a constanttemperature range of 25−105 °C at intervals of 20 °C.
min, the MM ethyl acetate solution (20 wt %) was added dropwise within 5 h. The polymerization was continued for additional 3−6 h after charging all MM. Upon completion, products were transferred to a second vessel and cooled to ambient temperature. Prior to characterization, polymerization products were precipitated with methanol and redissolved in ethyl acetate twice to remove the unreactive monomers and dried under vacuum at 30 °C for 40 h. Characterization of MMs and Graft Copolymers. 1H and 13C NMR measurements were made using a Bruker Advance 2B to determine monomer conversions and MM sequential structures, respectively. Molecular weights and molecular weight distributions of copolymers were characterized with gel permeation chromatography (GPC) using a Waters 1525/2414 GPC equipped with waters styragel HR columns and a differential refractive index detector. Measurements were made at a constant temperature of 30 °C using tetrahydrofuran as the eluent at a flow rate 1.0 mL/min. Reported values are relative to PS standards. Glass transition 6947
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ACS Omega Table 2. Characterization of Acrylic Copolymers Containing the Different Chain Length MMsa x (%)b run
comonomer
2-EHA
AA
comonomer
Mwc (kg/mol)
Đc
MwBd (kg/mol)
qe
Tg1f (°C)
1.83 4.25
92
14
−23.4 −47.8
5.84
Tg2f (°C)
PMA PL10C4
MA L10C4
100 96.5
∼100 ∼100
∼100 ∼100
135 139
PL20C8
L20C8
95.4
∼100
∼100
307
202
29
−29.4
PL50C20
L50C20
94.3
∼100
∼100
248
14.8
163
9
−50.3
−12.9
PL100C40
L100C40
92.7
∼100
∼100
313
18.8
205
8
−53.5
−4.5
a
b
1
Feed ratio of monomers: 2-EHA/comonomer/AA = 61:33:6 (wt/wt/wt). Conversion of monomers determined by H NMR analysis. Determined by GPC in THF and calibrated with PS standards. dMolecular weight of copolymer backbone. eNumber of branches in copolymer chain, q = (Mw,polymer × wMM)/Mw,MM, where wMM is the weight percentage of MM in the copolymer. fDetermined by DSC. c
Mechanical performance tests were carried out on films generated by dissolving polymers in ethyl acetate, coating the solutions on poly(ethylene terephthalate) films using a wirewound drawdown rod, and drying films under ambient conditions. Samples of each film were cut and weighed to confirm coating weights of 25 ± 3 g/m2 (≈1 mil). Tack force, peel strength, and shear holding times were measured according to ASTM standards, including ASTM D6195-03, ASTM D903-98, and ASTM D6463/D6463M-06. Both tack and peel testing were carried out with a Zwick/Roell Z020 Materials Testing Machine.
where nPLLA, nPCL, and ntotal represent the mole numbers of PLLA, PCL, and the total, respectively. Using eq 1, a value of 6600 g/mol is obtained, indicating that only PL100C40 may participate in chain entanglement. All MMs were found to possess a glass transition region and no melting transition when analyzed using DSC. Although homopolymers of both L-LA and ε-CL are crystalline, the random copolymer does not appear to form a crystalline phase at the molar ratio of 5:4. Table 1 shows a dependency between Tg and MM chain length with values increasing by more than 15 °C, moving from the smallest to the largest MM. It is expected that Tg is inversely related to Mn,45 and a plot of Tg values against the inverse of Mn,NMR values is fit accurately with the equation for a line (R2 > 0.97, Figure S4). As reported previously, L-LA has a greater reactivity than that of ε-CL when Sn(Oct)2 is used to catalyze the reaction.46,47 Thus, the reaction would be expected to produce copolymer chains composed of rich lactidyl units close to the MM head and terminated by a rich caproyl unit sequence. However, the presence of transesterification reactions results in a more random distribution of repeat units along the chain structure. Larger MMs require longer reaction times, which should provide for a greater influence on transesterification. Measurements made using 13C NMR show that the MMs L50C20 and L100C40 did tend to possess a more random distribution of lactidyl and caproyl units along the MM chains when compared to those of L20C8 and L10C4. However, the differences are small. The polyester branched acrylic copolymers were synthesized from 2-EHA, AA, and MM at a weight-based feed ratio of 61:6:33, as shown in Scheme 1. For comparison, an acrylic polymer is also produced by replacing MM with methyl MA at the same feed ratio. The properties for the polymers are summarized in Table 2. The different reaction runs to produce these polymers are indicated by the number of lactidyl units that compose the MMs. For example, the reaction run is indicated as PL10C4 when the L10C4 MM acts as the hard monomer to produce the polyester branch copolymer. With the exception of 2-EHA, monomers introduced, including MMs, were completely consumed in copolymerization reactions. Slightly lower conversion of 2-EHA is presumed to be due to its lower reactivity than that of methacrylic,47 and the conversion was further decreased with increasing MM chain length. From Table 2, it can be seen that the weightaverage molecular weight (Mw) of the copolymers increased when the MM was substituted for MA. Values for Mw as well as Đ increased as the MM chain length was extended. The broad molecular weight distributions in PL50C20 and PL100C40 are
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RESULTS AND DISCUSSION Preparation of MMs and Graft Copolymers. All MMs were generated with the same molar ratio of L-LA to ε-CL of 5:4 but with different HEMA contents, which produced a range of chain lengths, including the structures L10C4, L20C8, L50C20, and L100C40. The synthesis of MMs is shown in Scheme 1. The subscripts indicate the number of lactidyl (L) and caproyl (C) repeat units. Each L-LA produces two L repeat units in the MM chain, whereas ε-CL contributes a single C unit along the chain. The MM compositions, number-average molecular weights (Mn), and dispersities Đ are summarized in Table 1. Also listed are the Tg values and percentage of the MM chains possessing a reactive terminal double bond (Fvinyl) determined using 1H NMR. The 1H NMR spectra, GPC traces, and differential scanning calorimetry (DSC) thermograms of the MMs are shown in Figures S1−S3. The favorable comparisons between compositions determined with 1H NMR analysis and those based on feed ratios indicate that reaction conditions were sufficient to provide for complete conversion of monomers. The MMs produced with double bonds were nearly 100% for most structures but decreased slightly for the largest MMs. This is believed to be due to the longer reaction times required for the longest two MMs, which provided for some self-polymerization of double bonds. The Mn values determined using 1H NMR were nearly identical to target values and ranged from 1.42 to 12.47 kg/ mol. The Mn values determined using GPC were different from the theoretical values, which can be attributed to the use of PS standards as references. The GPC-determined Đ values were similar for all four MM samples, ranging from 1.65 to 1.76. The entanglement molecular weight Me for the polyester MMs can be estimated from values for PLLA (Me(PLLA)) and PCL (Me(PCL)) as41 Me = Me(PLLA) ×
n nPLLA + Me(PCL) × PCL ntotal ntotal
(1) 6948
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ACS Omega attributed to the difference in relative reactivities of acrylates and MMs.48,49 Their 1H NMR spectra, DSC thermograms, and GPC traces are shown in Figures S5−S7. Table 2 lists Tg values for the copolymers. The substitution of MA (Tg of 9 °C) with the L10C4 MM (Tg of −31.6 °C) resulted in a relatively large decrease in Tg from about −24 to −48 °C. For the much longer L20C8 MM (Tg of −24.8 °C), this gap shrank to less than 6 °C. For the two largest MMs, L50C20, and L100C40, the DSC thermal scans changed significantly with two Tg values being observed. The lower Tg value is thought to be associated with MM-poor resin phase and the higher transition with the polyester MM rich phase, which is attributed to a lower reactivity of the MMs with increasing chain length. The gap between Tg values expanded by more than 10 °C in moving from the L50C20 to the L100C40 MM. The results not only indicate a poor compatibility between the acrylic and lactide-based phases but also show that a critical chain length for the polyester MM is required to provide a great enough immiscibility to induce phase separation. This immiscibility then increased as the chain length of the MM was increased further. The presence of phase separation is likely an advantage for adhesive materials. Such polymers can reach different degrees of relaxation over a wide range of temperatures,50 and the presence of two phases provides more options to manipulate properties with additives. SPM images for films cast with the graft acrylic copolymers are shown in Figure S8. It is apparent that the polymer morphology changed with MM chain length, although some dots and parallel lines appear in the images owing to the use of tapping mode. The most evident distinction is between polymers PL10C4 and PL20C8 (Figure S8a,b) and polymers PL50C20 and PL100C40 (Figure S8c,d), with the latter pair with the longer chain MMs clearly demonstrating phase separation. The aggregation degree of each phase is not well-distributed due to the traditional free-radical copolymerization, leading to broad MM composition distribution in copolymers. Their phase images indicate the presence of a hard phase (light color) presumed to be aggregated polyester-rich regions surrounded by a softer matrix (dark color), which is likely the acrylic-rich phase. Influence of MM Chain Length on Linear Viscoelasticity. Shear rheology in the linear viscoelastic region of a polymer is quite sensitive to changes in microstructure.51−54 This approach is also used to characterize the branched acrylic copolymers. For the PL10C4 and PL20C8 polymers, normal stress changed rapidly for temperatures above 65 °C, resulting in the unreliable data. Thus, measurements for these samples were made below this temperature. The plots of complex viscosity (η*) versus frequency (f) for the different graft copolymers are shown in Figure S9. The master curves (reference temperature 25 °C) for the complex viscosities are shown in Figure 1. PL10C4 and PL20C8 start to show the transition from the Newtonian to non-Newtonian region under 10−2 Hz. However, the two samples possessing longer branches demonstrated non-Newtonian behavior at very low frequencies (10−6 Hz). Throughout the high-frequency region, the slopes of all samples are similar (i.e., −1/2) and close to that of the Rouse model in the intermediate frequency region, indicating that chain entanglement is not substantial. However, at low frequencies, the plot shows a clear distinction in behaviors for the copolymers due to the different MM chain lengths.
Figure 1. Master curve of complex viscosity of copolymers with different branch lengths generated using a reference temperature of 25 °C and the aT from the low-frequency fit of G′.
It has been found that molecular weight distribution and branch chain length affect the terminal frequency rheological properties and thus the zero shear viscosity, owing to the introduction of a long relaxation time component.30 The molecular weight distributions changed from 3.3 to 12.7 for linear low-density polyethylenes having similar molecular weight, resulting in the increase of zero shear viscosities, but the increment is within 1 order of magnitude.56 From Figure 1, the complex viscosity of 3.5 × 107 Pa·s can be observed for PL50C20 at aT × f of 2.8 × 10−6 Hz, where the Newtonian transition zone has not been reached yet. The fitting of the data using the Carreau−Yasuda55 equation provides a zero shear viscosity value for PL20C8 (3.8 × 105 Pa·s) that is about twice that of PL10C4 (1.7 × 105 Pa·s). Comparing the complex viscosity of 3.5 × 107 Pa·s for PL50C20 before reaching the Newtonian transition zone with the zero shear viscosity of 3.8 × 105 Pa·s for PL20C8, such a substantial viscosity difference cannot be attributed to the molecular weight distribution effect only. From the G′ master curves for the copolymer samples (Figure 2), it can be seen that at high frequency, all of the
Figure 2. Master curves of G′ for the copolymers with different branch lengths at the reference temperature of 25 °C. aT is determined from the superposition of G′ values at low frequency by horizontal shift.
curves showed a similar frequency dependency close to that expected from the bottlebrush Rouse model. The greatest G′ value at high frequency is for PL100C40. There is a great difference in the rheological behavior in the low-frequency region due to differences in phase separation and thus physical cross-linking. For the two shorter chain lengths, the increase of the molecular weight produces terminal relaxation time increases and the curve moves toward low frequency. Although the molecular weight of PL50C20 is lower than that of PL20C8, it still shows a longer relaxation time. Moreover, the slope of the low6949
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Figure 3. Master curves for copolymers of different branch lengths at the reference temperature of 25 °C in (a−c) and master curves for L100C40 at the reference temperature of 13.8 °C in (d). aT of PL50C20 and PL100C40 are determined from the superposition of G′ values at low frequency by horizontal shift. The solid markers indicate G′ values, whereas the open markers indicate G″ values.
from those for the shorter branch chain length polymers. Plots of the moduli overlap with a slope of 0.5, and it appears that the molecule chain is going to relax as with the bottlebrush Rouse model behavior at frequency lower than 10−6 Hz. However, the molecular weight of PL50C20 is lower than that of PL20C8 and the terminal relaxation of the bottlebrush molecular chain should appear at a comparable time scale as that of PL20C8 (∼10−2 Hz) even when PL50C20 has longer side chains. Therefore, the relaxation appears to be at least 104 times slower than the terminal relaxation of PL20C8 and cannot be ascribed to the molecular relaxation. In other words, there must be an extra elastic contribution to the G′ value.59,63 For the PL100C40 system, the elastic character (G′ > G″) at low frequency is obviously different from that demonstrated by the other samples. Since the branch length is about twice the Me, it is necessary to elucidate the role of arm relaxation. For the bottlebrush polymer having a side-chain length shorter than the backbone, the relaxation of side chains normally occurs before the relaxation of the backbone and a plateau at high frequency can be observed if entangled.60 Since the Tg of L100C40 is −15.7 °C while the higher Tg of PL100C40 is −4.5 °C, the reference temperature of the master curve for L100C40 was set at 13.8 °C when 25 °C is selected as the reference for PL100C40 (Figure 3d), which is to secure the comparison of both samples at the same friction state. Rouse-like relaxation behavior with terminal relaxation appears. Therefore, the relaxation of the whole bottlebrush molecule could be slightly slower than that of PL20C8 due to the difference in molecular weights. It indicated that the contribution of molecular relaxation was mainly covered by the contribution of phase separation. If there is phase separation, it would be expected that moduli from different temperatures cannot superpose well (Figure 3c,d). The failure of time−temperature superposition indicates existence of at least two different relaxation processes
frequency range from scale 2 of PL20C8 to scale 0.5 of PL50C20 indicates the existence of an additional relaxation mechanism, namely, the contribution of phase separation morphology. For the phase-separated samples, the slope of PL50C20 at low frequency is very close to 0.5, similar to the scale of block copolymer lamellar phase separation, whereas the scale of PL100C40 is 0.2, which is between the scale of columnar phase and globular phase.61,62 Therefore, the two samples with longer side chains were mainly influenced by phase separation. The log−log plots of G′ and G″ as a function of frequency for different graft copolymers are shown in Figure S10. The master curves for G′ and G″ are shown in Figure 3. Moduli from different temperatures can be superposed well. PL10C4 and PL20C8 displayed the terminal slopes of G′ and G″ as 2 and 1, respectively. At high frequency, moduli values nearly overlap with slopes of 0.5. The result is similar to the reported result on the Rouse model of bottlebrush systems.57 Although the Me of backbone (2-EHA) is about 37.4 kg/mol,58 the molecular weight of the shortest system, PL10C4, is about 2.45 times this, which means in theory, each system can entangle. However, the backbone entanglement is suppressed due to the introduction of branching, which increases its chain rigidity. Also, the average molecular weights between two branching points of the copolymers range from 6.6 to 25 kg/mol, which are lower than those of Me, and thus the copolymers are too small to provide for entanglement.32,60 Thus, PL10C4 and PL20C8 likely possess bottlebrush rheological behavior and the scaling exponent indicates no phase separation. Furthermore, the backbone of PL50C20 is shorter than that of PL20C8, whereas the arm length is much longer, indicating entanglement of the backbone of PL50C20 is not substantial. For PL100C40, although the backbone is longer than that of PL20C8, the much longer arm also suppresses backbone entanglement. When the branch length was extended to L50C20, the lowfrequency behaviors of G′ and G″ begin to change significantly 6950
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forms, which possessed high cohesive strengths, appears to provide an opportunity to increase tack while maintaining good properties overall. Thus, these graft copolymers provide a variety of means to tailor adhesive properties through MM and backbone composition, MM length structure, and use of additives. This last approach is likely to be greatly enhanced when MM chain lengths are sufficiently long, which promotes phase segregation. Given that PSA performance is typically governed to a greater extent by bulk viscoelastic properties rather than surface energy, great insights on the observed adhesive properties can be gleaned from the rheological data.12,64 The Dahlquist criterion was one the earliest connections between rheological data and performance.40 A more elaborate analysis is the Chang window, which identifies characteristic frequencies associated with various processes a PSA undergoes (e.g., tack, peel, and shear testing)65 and uses the dynamic moduli as axes to produce a performance map useful in product design.66 Chu provided an empirical connection between performance properties and rheology data based for resins, tackifying resins, and PSA.67 It is this analysis that provides useful insights on the graft copolymers studied here. Chu concluded that the G′ and Tg at the application temperature are the most important parameters in determining performance. For tape PSAs, a G′ range of 50−200 kPa and a Tg range of −15−10 °C is suggested. For label PSAs, ranges of 20−80 kPa and −30 to −10 °C are suggested for G′ and Tg, respectively. The use of Tg values is based on their assignment as the maximum in the tan δ temperature sweep. Temperature sweeps were carried out at a frequency of 10 rad/s or about 1.6 Hz. It is expected that the Tg values measured via DMA at this frequency will likely be slightly higher than those found with DSC. Table 3 provides G′ and G″ at 25 °C for a variety of
(molecular relaxation and domain relaxation) with different temperature dependencies. To illustrate in a quantitative manner, we could determine the shift factor aT from the superposition of G′ value under low frequency as aT(G′) and aT from superposition of G″ value under high frequency as aT(G″). Comparing the temperature dependence of aT of copolymers with different branching lengths under high and low frequencies (Figure 4), aT(G″) of PL50C20 and PL100C40 have
Figure 4. Relationship between aT and temperature of copolymers with different branch lengths for the reference temperature of 25 °C.
similar tendency to that of PL10C4 and PL20C8, indicating a similar molecular relaxation process at a high-frequency regime. However, aT(G′) curve of PL50C20 and PL100C40 obviously deviated from the curve of aT(G″). Such behavior clearly illustrates two different kinds of relaxation process due to the phase separation and chain relaxation, respectively. Effects of MM Chain Length on Mechanical Properties. The adhesive performance properties of the graft copolymer samples are shown in Figure 5. A change in mechanical properties from the conventional MA-containing copolymer to those generated with the MMs was evident. It can also be seen that tack, peel, and shear results were dependent on the chain lengths of MMs. This is an interesting result given that the copolymers contain the same MM content and composition overall. Although the copolymers were previously split into two distinct groupings, on the basis of phase separation, there was no clear distinction in performance between these groups. The sharp increase in peel and shear for PL100C40 is likely due to the large increase in chain length for this sample relative to that of the others. A promising finding is that tack did not decrease with increasing shear for PL50C20 and PL100C40. Both the peel and shear performance of the copolymers were quite good. However, the tack value would be considered a bit low by commercial adhesive standards. The composition of the MMs can be modified to enhance tack. However, this is at the expense of cohesive strength and results in a decrease of shear performance. The relatively small decrease in tack associated with moving from the shorter MMs to the longer
Table 3. Comparison of Storage Moduli at Different Frequencies for Graft Copolymers at 25 °C G′ (kPa)
G″ (kPa)
run
0.016 Hz
1.6 Hz
16 Hz
″ G16Hz ′ G0.016Hz
PL10C4
5.19
37.99
103
19.8
PL20C8
8.66
51.93
139
16.0
PL50C20
18.7
56.71
173
9.25
PL100C40
24.1
104.76
296
12.3
frequencies for the copolymers. It can be seen that for a frequency of 1.6 Hz, all G′ values for the copolymers were within the suggested range. For the short-chain MM copolymers, Tg values were slightly low, whereas for the longer MMs, two Tg values were obtained (see Table 2).
Figure 5. Performance properties including tack force, peel strength, and shear resistance for graft copolymers. 6951
DOI: 10.1021/acsomega.8b00737 ACS Omega 2018, 3, 6945−6954
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ACS Omega
polymers were strongly dependent on MM composition. The results presented here demonstrated that MM chain length provided another means to manipulate properties. The combination of these along with the use of additives for targeting-separated microphases provides a powerful set of tools for tailoring the performance of materials utilizing the graft copolymers.
Another empirical relationship suggested by Chu is based on the observation that a lower G′ is desired at low frequencies for better shear resistance, whereas a higher value is desired at higher frequencies to promote better peel results. As discussed earlier, G′ increases in a regular fashion, so Chu uses the value at 0.016 Hz (0.1 rad/s) as a gauge of its low-frequency value in combination with the ratio of G″ value at 16 Hz and G′ at 0.016 Hz as a gauge of how quickly values rise as the frequency is increased. From a large sample set, it was proposed that G″ PSAs with 20 kPa < G′0.016 Hz < 40 kPa and 5 < ′ 16Hz < 300
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ASSOCIATED CONTENT
S Supporting Information *
G0.016Hz
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsomega.8b00737.
possess a good balance of performance properties. It is interesting that only PL100C40 meets these criteria. It is likely that the long side chain is directly responsible for this result, which bolsters the case that the high-biomass content copolymers that were found to demonstrate good adhesive properties can be further enhanced through the graft length and phase separation.
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CONCLUSIONS The work reviewed above clearly shows that polyester branch chain length governed a large part of the properties of the generated hybrid acrylic polymers. Thermal analysis identified a split of a single glass transition into two second-order transitions when the L50C20 or L100C20 MM replaced L20C8 or L10C4 at the same mass fraction in the synthesis. The thermal separation between the transitions grew significantly larger when L100C40 was used as the MM. The behavior indicates the likely formation of acrylic and polyester phases, which were demonstrated visually in SPM phase images. Rheological analysis of the linear viscoelastic responses was also used through small-amplitude oscillatory shear, and dynamic master curves were constructed by time−temperature superposition. The data indicate that all graft copolymers underwent the bottlebrush Rouse-like behavior and the copolymers containing the L50C20 or L100C40 MM clearly demonstrated phase separation. It is shown that with the change of the side-chain length, different types of phase separation occur and enhance the mechanical properties of the systems. The effect on adhesive performance is expected to be twofold. First, phase separation produced a low Tg MM-poor phase and a high Tg polyester-rich phase. The combination provided a soft polymer that could retain cohesive strength. This is believed to be exemplified by the retention of tack for the copolymer containing the L100C40 MM, which demonstrated the highest shear values. Furthermore, the separation into immiscible microphases provides a means to manipulate transitions and properties separately, allowing the performance of such species to be further modified with chemical additives (e.g., tackifying resin). Second, the presence of relatively long branches and phase separation are believed to enhance the cohesive strength of the copolymer, which increased both its peel and shear performance. The results of performance testing were consistent with this assertion. The most evident connection is that between the elastic contribution of phase separation observed for the copolymer containing the L100C40 MM and its measured shear resistance. The time reported is that for the adhesive to flow from the surface resulting in adhesive bond failure. Increasing side-chain length will extend the time required for such processes. It should be noted here that no effort is made to optimize the performance of the polymers. Previously, it is demonstrated that physical and performance properties of such
All 1H NMR spectra, GPC plots, Tgs of MMs and copolymers, SPM images, plots of complex viscosity, G′ and G″ versus frequency for copolymers with different side-chain lengths (PDF)
AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected] (W.Y.). *E-mail:
[email protected] (S.J.S.). *E-mail:
[email protected] (W.-J.W.). ORCID
Wen-Jun Wang: 0000-0002-9740-2924 Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The authors acknowledge National Key Research and Development Program of China (2016YFB0302400), National Natural Science Foundation of China (21533611, 21420102008, and U1462115), and Chinese State Key Laboratory of Chemical Engineering at Zhejiang University (SKL-ChE-15D03) for financial support.
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REFERENCES
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