PTHF Segmented Block

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

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Bimodal Crystallization Kinetics of PBT/PTHF Segmented Block Copolymers: Impact of the Chain Rigidity Andre ́ de Almeida,* Matthias Neb́ ouy, and Guilhem P. Baeza Univ Lyon, INSA-Lyon, CNRS, MATEIS, UMR5510, F-69621, Villeurbanne, France

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S Supporting Information *

ABSTRACT: By combining linear rheology and differential scanning calorimetry experiments performed under isothermal and nonisothermal conditions, we clarify the mechanisms of crystallization occurring in three industrially relevant PBT/PTHF segmented block copolymers. After a careful and systematic analysis of the crystallization kinetics based on “classic” Avrami-like models (1 < n < 2), we reveal (and quantify) the twostep nature of the PBT segments’ association. Besides, we discuss the role of the hard (PBT) and soft (PTHF) segments’ length by confronting our results to PBT homopolymers. As expected, while 2 kg mol−1 PTHF segments are found to delay substantially the crystallization of the PBT within the copolymers, we demonstrate that shortening them down to 1 kg mol−1 results in similar kinetics as for neat PBT lamellar crystallization. This result, raising fundamental questions on the chains’ conformation, is finally discussed jointly with our rheological tests (gelation) and recent theoretical predictions with the aim to bring new insights into the topology of such complex systems.

1. INTRODUCTION Thermoplastic elastomers (TPEs) made of segmented (or “multiblock”) copolymers consisting of a sequence of hard and soft segments (HSs/SSs) are widely used in the industry for their polyvalence and ease of processing.1 They notably occupied for decades a good place in the automotive sector as well as in various fields such as insulating sheaths,2 bitumen modifiers,3 and adhesives.4 This great popularity mostly comes from their versatile mechanical properties originating from the phase separation of their constituents, eventually tunable by modifying, e.g., the architecture or the chemistry of the copolymer. As a matter of fact, while at the service temperature, “free” SSs give the material its flexibility; HSs tend to aggregate into amorphous5,6 and/or crystalline7,8 domains leading to the formation of a rigid network. Based on the relative fraction of those two phases, different versions of the same TPE may have in consequence the possibility to operate as either energy dissipators9 or reinforcing fibers.10 Besides, the recent synthesis of bio-based11 and biodegradable12 TPEs together with their ability to be reshaped and potentially recycled (contrary to classic vulcanized rubbers) makes them promising candidates for the development of innovative green technologies. Among TPEs, poly(ether−ester)s consisting of a sequence of poly(butylene terephtalate) (PBT) and poly(tetramethylene oxide) (or polytetrahydrofuran, PTHF) units are synthesized industrially for more than 40 years for their outstanding mechanical behavior, chemical resistance, thermal stability, and rapid crystallization.13−18 However, in spite of an intense research activity aimed at understanding their structure− properties relationship, a clear picture of the network formation along with the HSs crystallization still remains © XXXX American Chemical Society

elusive. In fact, a great part of the work published in the TPEs field consists of describing, in a parallel way, the impact of the chain composition on (i) the macroscopic mechanical properties and (ii) the microstructure of hard domains. In spite of its undeniable usefulness in terms of empirical understanding, this method has however not succeeded yet in describing that complex relationship. To further extend our comprehension, we thus believe that a deeper theoretical approach completed with molecular dynamics simulations is necessary. In this logic, we proposed recently an original and quantitative model19 based on SAXS measurements20 allowing to rationalize the increase of the storage modulus with the HSs fraction in well-defined TPEs. In the same vein, it is worth to remind that a wide variety of experiments21 and simulations22 performed on semicrystalline homopolymers have pointed the importance of the polymers’ topology. The role of the tie molecules has notably been at the center of the attention,23 being of great interest to inspire similar concepts in block copolymers. Beyond probing the static and dynamic behaviors of TPEs, both their crystallization and gelation kinetics have been extensively studied to extract structural information at various length scales. The former has indeed been the object of many articles among which the empirical Avrami model24,25 and its derived forms26−28 describing the crystallization from a characteristic time (or speed “K”) and a growth mode (or geometry “n”) have emerged as references. The latter has been investigated in a more rational way by Winter and co-workers, Received: August 6, 2018 Revised: January 11, 2019

A

DOI: 10.1021/acs.macromol.8b01689 Macromolecules XXXX, XXX, XXX−XXX

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Figure 1. (a) Chain morphology of HS30, HS40, and HS65. PBT and PTHF are respectively represented in blue and red. (b, c) PTHF and PBT repetition units.

Table 1. Microstructure of TPEs sample

HS (wt %)

⟨N⟩

MHS (kg mol−1)

MSS (kg mol−1)

M̅ n (kg mol−1)

Ip

Tma (°C)

Tca (°C)

T0mb (°C)

HS30 HS40 HS65

30 40 65

7.7 6.5 7.7

1.0 1.5 2.1

2.0 2.0 1.0

25 25 25

2 2 2

180 200 208

96 136 169

187 205 212

a

Data extracted from the DSC thermograms presented in Figure 2. bData extrapolated from the Hoffman−Weeks equation (see Supporting Information section 2.1). “PTHF”) of various molecular weights (1 and 2 kg mol−1 in the present work). Tetrabutyl titanate (TBT) is used as a catalyst, and the polyether is protected against oxidation during polymerization by 1,3,5-trimethyl-2,4,6-tris (3,5-di-tert-butyl-4-hydroxybenzyl)benzene. The polycondensation is performed at a temperature of 250 °C under vacuum, with a typical polymerization time of 200 min. Additional details on the chemical route can be found in ref 18. Subsequent to the polymerization, the material is extruded as strands and pelletized. The so-formed granules are subsequently dried at 100 and 140 °C for 2 and 5 h, respectively, and used as such in the following study. For the sake of clarity, we report in both Figure 1 and Table 1 the details on the chains’ microstructures. The latter also contains the thermal characteristics of the three TPEs, i.e., the crystallization and melting temperatures (Tc and Tm) as well as the Hoffman−Weeks extrapolation (T0m, see Supporting Information section 2.1 for details). The corresponding thermograms are presented in Figure 2 where the samples were isothermally crystallized at 140 °C (HS30), 170 °C (HS40), and 190 °C (HS65) for an hour prior to the measurement.

who discussed the network’s morphology close to the gel point,29,30 being defined rheologically as the time required by the loss factor (tan(δ)) to exhibit a flat behavior as a function of the frequency. In this framework, a relevant example of such kinetics investigation lies on the network growth in neat PBT (which crystallizes in a triclinic structure).31 While short chains (23 kg mol−1) characterized by a high mobility at long-range were observed to crystallize quickly in a spherulitic way (n > 3 for neat PBT and PBT mixed to montmorillonite),32 longer ones (115 kg mol−1) were restricted to a much slower formation of lamellae (n < 2 for neat PBT and PBT blended to epoxy).33 Similar trends were observed in TPEs where the HSs association seemed to be delayed for longer copolymers,34,35 but also longer or more polydisperse HSs,36,37 as well as softer chains (i.e., lower HS content),38 all these parameters further lowering the crystallites’ content.7 The article is organized as follows: After a quick description of the copolymers’ microstructure, we provide all the experimental details in section 2. We then present successively the results coming from isothermal rheology and calorimetry in sections 3.1 and 3.2. We further complete our study with postcrystallization temperature ramp tests that we show in section 3.3. Finally, we discuss our results from both experimental and theoretical point of views in section 4.

2. MATERIALS AND METHODS 2.1. Materials. The copoly(ether−ester)s investigated in this article (HS30, HS40, and HS65) are industrially relevant TPEs synthesized by DSM, The Netherlands. They are made of a sequence of hard (poly(butylene terephtalate), PBT) and soft (polytetrahydrofuran, PTHF) segments and contain respectively 30, 40, and 65 wt % in PBT. While their total molecular weight is fixed to M̅ n = 25 kg mol−1 with a polymolecularity index Ip = 2 (Flory statistics), the SSs they contain are either MSS= 2 or 1 kg mol−1 respectively for HS30 and HS40 on the one hand and HS65 on the other. The average molecular weight of the HSs M̅ HS and their number along the chains ⟨N⟩ are calculated from their mass fraction (wt % HS) and the size of the SSs used as “separators”. These materials are produced via melt transesterification and subsequently polycondensation of dimethyl terephthalate (DMT), 1,4-butanediol (BDO), and poly(tetramethylene oxide) (“PTMO” or

Figure 2. DSC thermograms of the three TPEs crystallized isothermally at 140 °C (HS30), 170 °C (HS40), and 190 °C (HS65). The samples are first cooled at −90 °C and then heated to 250 °C at 10 °C min−1 (“melting”), maintained at 250 °C for 5 min, and finally cooled to −90 °C (“crystallization”) at the same rate. B


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Figure 3. Storage modulus as a function of time G′(t) for (a) HS30, (b) HS40, and (c) HS65. Up to ca. 100 s, the temperature is decreasing from T > Tm down to Tiso before stabilizing. Black solid lines are fit to the data (see Supporting Information section 1.3 for details). (d) Corresponding values of G′inf as a function of ΔT. Dashed lines are linear fits to the data.

Figure 4. (a) Evolution of the storage (G′, diamonds) and the loss (G″, solid line) moduli with time for HS30 crystallized at Tiso = 140 °C and ω = 10 rad s−1. Colored arrows are related to Figure 9. (b) Time required to reach the gelation point in HS30, HS40, and HS65 measured at different Tiso. min−1) at Tiso and maintained to that temperature for crystallization. The time required to reach the equilibrium was estimated from rheological experiments, corresponding to a saturation of the storage modulus G′(t) = G′inf (see Figure 3a). (Also, we systematically checked that no exothermal event was occurring after this time.) The heat flow was recorded during the isothermal crystallization and used to quantify the crystallization kinetics (see Figure 5a). Once the crystallization completed, the samples were heated from Tiso to 250 °C at a fixed heating rate of 10 °C min−1 to melt the soformed crystallites, providing additional information about their morphology (see Figures 8 and 9a). The corresponding crystallinity ratio (i.e., the mass fraction in crystallites) was systematically calculated from Xinf = ΔHm/ΔH*, where ΔHm is the experimental

Both Tc and Tm are found to be similar to the values reported in our previous work on nonisothermally crystallized samples.39 2.2. Sample Preparation. The 1.0 mm thick films were prepared through hot-pressing (2 bar) of TPE granules in a Teflon mold. The temperature was maintained for 9 min at 220, 240, and 250 °C for HS30, HS40, and HS65 respectively, i.e., roughly 40 °C above the melting point of the PBT crystallites (see Table 1). The samples were then removed from the hot press and cooled at ambient temperature in air. 2.3. Methods. Differential Scanning Calorimetry (DSC). DSC experiments were performed on a Q8500 (PerkinElmer, USA) apparatus under a nitrogen flow of 20 mL min−1. Hot-pressed samples of about 10−15 mg were first heated to 250 °C for 5 min to ensure complete melting. They were subsequently cooled quickly (50 °C C


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Figure 5. (a) DSC crystallization exotherms for HS65 at Tiso = 175 and 180 °C; the dashed line is a fit to the data with a fourth-order polynomial. (b) DSC thermograms obtained from HS65 showing the crystallization under isothermal conditions for Tiso = 170, 175, 180, and 190 °C. Endothermic heat flow up. melting enthalpy and ΔH* = 145.5 J g−1 is the theoretical enthalpy of a perfect PBT crystal.18 Rheology. Rheological measurements were all performed in a strain-controlled rheometer (ARES 2kFRTN1 from Rheometric Scientific, currently TA, USA) using Invar parallel plates of 10 mm diameter (1.0 mm in sample thickness). A nitrogen convection oven allowed maintaining inert atmosphere with a temperature control better than ±1 °C. Time sweep experiments (Figure 3) consisted of heating the samples above their melting point and cooling them to Tiso at ca. 40−50 °C min−1 (see Supporting Information section 1.1 for temperature profiles). Typical strain amplitude and frequency were respectively set to 3−5% (linear) and 10 rad s−1. The gap was adjusted along the experiment to avoid strong (negative) normal forces caused by the sample contraction during the crystallization. In addition, dynamic temperature sweep experiments (Figure 9b) were performed on the HS30 sample right after its isothermal crystallization for 4, 17, 25, and 70 min spent at 140 °C.

Supporting Information section 1.2 frequency sweep experiments performed prior and subsequently to the crystallization procedure for each TPE at a given Tiso, evidencing further the gelation mechanism. Remarkably, a closer look at the data reveals a double-S shape for HS40 regardless of Tiso. The latter, actually discernible in HS65, too, suggests that the elastic network made of soft-strand-bridged crystallites19 is built through two successive steps, most likely related to the concomitant crystallites’ growth and soft strands’ rearrangement. This scenario is further supported by a series of semiempirical fits that we performed on the HS40 and HS65 data sets by using a two-mode exponential rise function reminding the well-known Avrami model (see Figure 3 and Supporting Information section 1.3). Note that in spite of the apparent single S shape of the HS30 data, we are convinced of the presence of a twomode building process similarly as in HS40 and HS65 as evidenced later in the article (see Figures 8a and 9a). Once the gelation was undoubtedly identified, we were then interested in quantifying its characteristic time tgel, often chosen as the G′ − G″ crossover position40,41 (see Figure 4a). The more robust definition of tgel proposed by Winter et al. lying on the independence of the loss factor tan(δ) = G″/G′ with the frequency40−42 is compared to the crossover method in Supporting Information section 1.4 showing a remarkable agreement. On the basis of this encouraging result and on the fact that multiwave experiments were hardly conductible over the large range of temperatures and chemical structures treated in this article, we decided to define the gelation as the G′ − G″ crossover position in the following. As one may have expected from classic thermodynamics,40 tgel is found to decrease exponentially with ΔT for the three TPEs (see Figure 4b), following eq 1.

3. RESULTS 3.1. Linear Rheology along Crystallization. To probe the crystallites’ network building along with time, we performed a series of time sweep experiments in the linear regime starting from T > Tm down to Tiso, at which the samples were isothermally crystallized. In Figure 3a−c, we present the evolution of the storage modulus as a function of time G′(t), at different Tiso, for the three TPEs. HS30, HS40, and HS65 are respectively crystallized between 100−155, 135−170, and 170−195 °C, corresponding to ΔT = T0m − Tiso ranging from 20 to 90 °C (Figure 3d). The temperature is stabilized at Tiso after ca. 4 min explaining the trend of G′(t) between 10 and 100 s (Supporting Information section 1.1), corresponding in reality to a temperature ramp. Following this well-controlled thermal history, the modulus exhibits then systematically a sigmoid-like shape, unambiguously related to the PBT segments association, i.e., with the formation of an elastic network spanning the whole material. In fact, Figure 4a evidences the typical transition from liquidlike (G′ < G″) to solidlike (G′ > G″) behavior after a few minutes spent at Tiso. Besides, while increasing ΔT is found to accelerate the crystallization in an expected manner, it further provides the TPEs with a (linearly) growing steady-state storage modulus (G′inf) as shown in Figure 3d. In particular, the ratio G′inf(Tiso‑min)/G′inf(Tiso‑max) is found to be close to 10 for the three TPEs, suggesting that Tiso impacts significantly the amount of crystallites at equilibrium. Indeed, despite a higher temperature, the sole soft phase can hardly be thought of being responsible for 1 order of magnitude loss on the storage modulus. For the sake of completeness, we provide in

tgel(t ) = t0 exp( −bΔT )

(1)

Here, t0 is the extrapolation of the gelation time at the melting point, and b is a coefficient inversely proportional to some activation energy, normalized by the perfect gas constant R. These parameters are reported in Supporting Information section 2.2 together with the characteristics on the crystallization kinetics measured from DSC. The evolution of tgel as a function of the temperature and the chain composition is discussed together with the crystallization kinetics in section 4. 3.2. Crystallization Kinetics Probed through Isothermal DSC. Motivated by the reproducibility and the apparent D


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Figure 6. Relative crystallinity X(t)/Xinf as a function of time at different Tiso for (a) HS30, (b) HS40, and (c) HS65. (d) Overall crystallization rate, t0.5−1, for HS30, HS40, and HS65 as a function of ΔT.

exhibiting flatter profiles (see Figure 5b), at first sight in concordance with the results from our rheological study. For most of the experiments performed on the three TPEs, point B was found at the same heat flow level as point A (see e.g. Tiso = 180 °C in Figure 5a), allowing thus a straightforward determination of the crystallization ratio through eq 2. However, for the lowest values of Tiso (see e.g. Tiso = 175 °C), point B was found significantly higher than point A, resulting in an uncompleted exotherm. To eliminate these experimental limitations, unambiguously related to the time needed by the sample to reach a thermal equilibrium, we decided to fit the exotherms at short time with 4-order polynomials (dashed line) and extrapolated them to “negative” times. The latter procedure, schematized in Figure 5a, allowed to evaluate the actual position of A, opening the way to a reliable determination of X(t) regardless of Tiso. On the basis of the above-mentioned procedure and in analogy with Figure 3, we present in Figure 6a−c the evolution of the crystallinity ratio along with time for the three TPEs at various Tiso. As one may have anticipated from the shape of the endotherms in Figure 5, X(t) always exhibits a sigmoidal shape with faster crystallization kinetics observed when ΔT is increased, i.e., with decreasing Tiso (see Figure 6). This usual trend has been reported many times for PBT-based polymers,32,33,43,46,47 explaining qualitatively the raise of the modulus along with time observed in Figure 3. In a more quantitative way, we present in Figure 6d the thermal dependence of the overall crystallization rate through the reciprocal crystallization half-time t0.5−1 (defined as the inverse of the time required to complete 50% of the relative crystallinity) as a function of ΔT. The latter results are found

robustness of our rheological experiments, we propose in this section to study the crystallization kinetics through isothermal calorimetry experiments with keeping the values of Tiso unchanged with respect to the previous section. The parallel between rheology and DSC will be then discussed in section 4. Direct correlation between the characteristic times is also provided in Supporting Information section 3. Overall Crystallization Kinetics. In Figure 5a, we present the exothermic transition dH/dt(t) corresponding to the crystallization of the sample HS65 under isothermal conditions for Tiso = 175 and 180 °C. Corresponding data for HS30 and HS40 are provided in Supporting Information section 2.3. The relative degree of crystallization along time, X(t), was calculated from the area under the exotherm from point A (assumed to be the crystallization onset) up to a given time. The latter was subsequently normalized by the total exotherm area (from A to B)38,43−45 as expressed in eq 2 in which dH/dt is the heat flow rate: t

X (t ) =

∫A ( ddHt ) dt B

∫A

( ddHt ) dt

(2)

Here it is worth mentioning that point A is found to be Tiso dependent and is preceded by a short-time period necessary for the temperature equilibration. As expected, the exothermic heat flow is first increasing from A until reaching a maximum (synonymous of a maximum in crystallization rate), before falling and becoming nil at B, marking the end of the crystallization process. At higher Tiso, the crystallization maximum is delayed to longer times with the exotherms E


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Figure 7. Crystallization kinetics analysis of HS65. (a) Overall crystallization, (b) primary crystallization, and (c) secondary crystallization. Solid and dashed lines are fit to the data based in eqs 5 and 6, respectively. (d) Primary and secondary crystallization rates as a function of ΔT. Solid and dashed lines are fit to the data (see the definition of t0.5−1 in the text) respectively standing for the primary and secondary crystallization processes.

to fit well with a single-mode exponential function as already shown in the literature for both TPEs43,44,48 and neat PBT.32,49 More interestingly, while almost identical t0.5−1 (ΔT) values are found for HS30 and HS40, much faster crystallization kinetics are observed for HS65. The latter finding suggests therefore that the molecular weight of the SS, rather than the HS fraction (or length), is the critical parameter controlling the phase transitions kinetics and, by extension, the topology of the network. Modeling Complex Crystallization Kinetics. The expression of the relative degree of crystallization during isothermal processes is generally described by the Avrami model (eq 3):

in which n and log(K) are respectively the slope and the intercept. In Figure 7a, we provide the corresponding representation for the sample HS65 at different Tiso (similar data for HS30 and HS40 are presented in Supporting Information section 2.4). Interestingly, after a first linear regime observed at short times, the DSC data systematically show a break in slope at later stages (particularly at low Tiso), suggesting once again the presence of different mechanisms along the crystallization and limiting therefore the use of eq 4 to a restricted time range. This deviation from the theoretical Avrami model has been observed in the literature for many different semicrystalline polymers26−28,38,43,51 and has been treated according to two different logics. While in some cases32,33,38 the authors simply mention that the Avrami model is well adapted for short times only, without going further into a possible physical explanation, other studies aim at developing more complex models providing a quantification of the crystallization during the whole process.26−28 In particular, the model proposed by Vehroyen et al. consisting of two successive Avrami-type crystallization events seems well adapted in our case.28 It is formalized through eqs 5 and 6 where k and k′ are respectively the rates of the primary and the secondary crystallization, n and n′ are the corresponding nucleation−growth modes, and tp,end is the time for which the crystallization switches from the primary to the secondary mode (i.e., the position of the “break in slope” in Figure 7a).

t

X (t ) =

∫A ( ddHt ) dt B

∫A

dH dt

( ) dt

= 1 − exp( −Kt n) (3)

where X(t) is the fraction of crystal phase as a function of time, and n and K are Avrami constants giving information about the crystallization mechanisms involved. While the former provides information about the nucleation type (sporadic or predetermined) and the crystal growth geometry (e.g., spheres, rods, or lamellae), the latter informs about the crystallization rate constant.26,50 These parameters are usually determined by plotting the double-logarithm linearization of eq 3 leading to log(− ln(1 − X (t ))) = log(K ) + n log(t )

(4) F


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Macromolecules Table 2. Avrami Parameters Calculated for the Primary and the Secondary Crystallization overall crystallization

primary crystallization

secondary crystallization

sample

Tiso (°C)

ΔTa (°C)

t0.5−1 (min−1)

n

k (min−n)

t0.5−1 (min−1)

n′

k′ (min−n′)

t′0.5−1 (min−1)

HS30

110 120 130 140 150 150 160 170 175 180 190

77 67 57 47 37 55 45 42 37 32 22

0.33 0.69 1.63 6.48 10.33 2.87 13.49 2.32 3.64 5.95 18.11

1.32 1.3 1.52 1.3 1.41 1.34 1.36 1.88 1.79 1.9 1.77

1.69 1.08 0.72 0.38 0.26 0.33 0.2 0.61 0.36 0.22 0.099

0.32 0.49 0.64 1.4 1.91 1.58 2.5 0.6 1.06 1.62 3.93

1.17 0.87 1.08 1.18 1.17 1.16 1.15 1.08 1.12 1.21 1.16

4.84 1.78 0.8 0.13 0.087 0.26 0.091 0.8 0.39 0.22 0.076

0.12 0.45 0.8 4.4 6.81 2.3 6.62 0.8 1.58 2.58 7.86

HS40 HS65

ΔT = T0m − Tiso.

a

Figure 8. DSC thermograms recorded subsequently to the isothermal crystallization procedure. The temperature is increased from Tiso up to 250 °C at 10 °C min−1, showing three endotherms at TII, TI, and TIII for (a) HS30, (b) HS40, and (c) HS65. Endothermic heat flow up. (d) Total weight fraction in crystallites (determined from the TII and TI melting peaks) as a function of Tiso for the three TPEs. Solid lines stand for linear fits.

X p(t ) =

X s (t ) =

∫0

t

t

( ddHt ) dt ( ddHt ) dt

∫0 p,end

= 1 − exp( −kt ) (5)

t − t p,end

( ddHt ) dt ∞ ∫t ( ddHt ) dt

∫t

lines are fit to the data with eqs 5 and 6, respectively (same mathematical expression for different time range). In Table 2, we report all the parameters, i.e., n, n′, k, and k′ for the three TPEs in a wide range of Tiso. Also, we provide the values of t0.5−1 and t′0.5−1 that one can calculate in a straightforward way from the latter parameters through t0.5−1 = (ln 2/k)1/n and t′0.5−1 = (ln 2/k′)1/n′ (see also Figure 7d). Remarkably, the n values are never integers and appear to be almost independent of Tiso (or ΔT) regardless of the TPE formulation. However, n is significantly higher in HS65 (mean value of 1.83) in comparison with HS40 or HS30, for which similar values are found (1.35 and 1.37, respectively), in

n

p,end

= 1 − exp( −k′(t − t p,end)n′)

p,end

(6)

The analysis of the primary and secondary crystallization kinetics for HS65 is presented in Figure 7b,c where straight G


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heated to 250 °C at 10 °C min−1 to probe (indirectly) the structural characteristics of the so-formed objects. At first sight, three endothermic peaks are observed regardless of the TPE composition (see Figure 8c as an example): (i) a minor transition at TII (secondary crystallization), for which the position is found to evolve steadily with the isothermal crystallization temperature Tiso; (ii) a large-amplitude melting endotherm at TI (primary crystallization) moving according to the Hoffman−Weeks plots provided in Supporting Information section 2.1;48,58 and (iii) a modest endotherm situated at a higher temperature TIII for which the amplitude seems to increase with lowering Tiso. In addition, the sample HS65 exhibits an extra peak located between TII and TI for Tiso = 170 and 175 °C. This event is sometimes observed in the literature and is believed to originate from the same physical mechanism as (ii).43,45 Because it is rather weak and close to TI, it will not be treated specifically in the following. On the basis of these results and in agreement with the literature on polypropylene45 and on neat PBT,59,60 it is clear that the major endotherm located at TI is related to the melting of the primary crystallites formed at early stages of the isothermal crystallization. In addition, we believe that the minor endotherm appearing at TII stands for the melting of secondary crystallites, likely to come from the crystallization of constrained HSs. Finally, we tentatively attribute the endotherm observed at TIII to the melting of few crystallites eventually re-formed during the heating scan as suggested in refs 45, 59, and 60. It is worth mentioning that the mechanism and extent of recrystallization process during the heating scan depend greatly on the stability of the primary and secondary crystallites formed at Tiso and of the heating rate used for the melting. More quantitatively, while the melting temperature TI is observed at around 180, 195, and 210 °C respectively for HS30, HS40, and HS65, it was reported close to 221 °C (significantly higher) in spherulitic PBT (Celanex 2000).60 In the same way, the total weight fraction in crystallites (Xinf) in TPEs, extracted from the TI and TII peaks, appears lower than the corresponding value in that neat PBT. As a matter of fact, the maximal values of Xinf reported for HS30, HS40, and HS65 are respectively 7.5, 15, and 23 wt % (see Figure 8d), clearly lower than the 47 wt % reported for the homopolymer.32 As one may have anticipated, it therefore appears that both the melting temperature TI and the crystallite content Xinf increase monotonically with the fraction of PBT in the chains. That trend may be understood by seeing the SSs as structural defects, limiting the growth of the crystallites from a quantitative point of view (fewer HSs can associate) as well as from a qualitative one (no spherulitic, i.e., 3D growth is allowed).61 Overall, all the above-mentioned arguments most likely explain the experimental results showing that the effective fraction of crystallized PBT units varies roughly between 1/4 in HS30 and 1/3 in HS40 and HS65. Beyond that static analysis, the quasi-fixed position of TI further indicates the unchanged size of the largest crystallites with changing Tiso for the three TPEs. On the contrary, Tiso is found to have a strong impact on the secondary structures, melting systematically right above their temperature of formation (Tiso). To further investigate the metastable nature of these crystallites, we propose below to perform similar experiments with the particularity to stop the crystallization at different stages, allowing to probe the successive states of the network building. The results are presented in Figure 9a where

agreement with the overall crystallization in Figure 6d. Because n falls systematically into the same range (between 1 and 2), the nucleation and growth mode of the PBT primary crystallization is assumed to be similar in the three copolymers. Indeed, according to the Avrami model, these values indicate a sporadic nucleation accompanied by a lamellar growth mode26,50 as further evidenced in our previous study from both SAXS and AFM experiments.39 Also, this type of structure has already been observed in the literature for comparable systems based on PTHF soft segments,52,53 reporting notably n = 1.5 (Zhu et al.).16 In the latter article, the composition of the TPE was close to the HS65 sample, made of 72% of PBT HS of ca. 3 kg mol−1 connected through PTHF units of 1 kg mol−1, suggesting somehow a universal behavior for this kind of copolymer. One can finally remark that the n values reported in Table 2 are much smaller than the ones determined for neat PBT (23.2 kg mol−1) crystallizing in a spherulitic way. In this case, n was found to vary between 3.83 and 4.15 for 197 < Tiso < 201 °C32,54 (see also section 4). Besides, noninteger n′ values relative to the secondary crystallization are also found to be independent of Tiso but are significantly smaller than their n counterparts. Remarkably, they all vary around 1.1 regardless of the TPE composition, suggesting a similar secondary crystallization mechanism in the whole sample set. Here again, similar values were reported in the literature for neat PTT and PET showing respectively 0.8 < n′ < 1.1 for 177 < Tiso < 207 °C47 and 1 < n′ < 1.3 for 211 < Tiso < 221.5 °C.43 This quasi-1D crystal growth (n′ ≈ 1) is generally attributed to lamellae thickening, crystallites impingement, crystals perfection (crystallization between the lamellae), or lamellae stacks when the primary crystallization occurred in a spherulitic growth mode.47,55,56 Nevertheless, for different growth modes, it can also correspond to the development of needles of finite dimensions.50 We believe that the secondary process is, in our case, related to the formation of small ribbon pieces or lamellar-like objects (also called “baby ribbons”57) subsequent to the gelation, limiting in consequence a good arrangement of the HSs. Moreover, one can obtain additional details on the temporal aspects of the crystallization processes by analyzing t0.5−1 and t′0.5−1. Here, these two parameters are preferred to k and k′ because of their physical meaning. Indeed, while the two formers are expressed in min−1 synonymous to “crystallization rate”, the two latter’s unit is min−n, making thus their interpretation intricate, in particular because of the variations of n. As clearly evidenced in Figure 7d, both the crystallization rates for primary and secondary crystallization increase exponentially with ΔT in a similar way as the overall crystallization process in Figure 6d. Also, this representation reveals that in spite of their different chemical composition, HS30 and HS40 are characterized by identical crystallization kinetics whereas the crystal growth seems to happen much faster in HS65. These data, which complete the qualitative evolution of n, further support the rheological study developed in section 3.1 in a quantitative way. It finally corroborates the two-step nature of the network building for which we propose a deeper characterization through a nonisothermal procedure in the next section. 3.3. Subsequent Melting of the Crystallite Network Built at Tiso. In Figure 8a−c, we present temperature ramp DSC experiments showing the melting of the crystallites network built at Tiso for the three TPEs. Once the crystallization completed (i.e., t = tinf), the samples were H


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Figure 9. (a) DSC thermograms recorded from HS30 samples heated to 250 °C at 10 °C min−1 after isothermal crystallization at Tiso = 140 °C for 4, 17, 25, and 70 min. Endothermic heat flow up. (b) Dynamic temperature sweeps (10 °C min−1) performed on HS30 samples after an identical isothermal crystallization performed in the rheometer (10 rad s−1, 5% strain). Inset: derivative of the storage modulus with respect to the temperature highlighting the breaks in slope observed at TII and TI.

to the behavior of much longer PBT chains (115 kg mol−1) exhibiting a lamellar growth.33 Hence, to compare in a suitable way these different data sets, we built master curves describing the relative crystallinity as a function of the normalized time, t/ta, on the basis of the work by Winter40 and Nojima.34 In this context, the overall relative crystallinity becomes independent of the temperature and can be expressed as

fresh HS30 samples were crystallized for 4, 17, 25, and 70 min at Tiso = 140 °C (see the arrows in Figure 4a for the corresponding time sweep curves). Strikingly, while a first and massive peak grows from the early stages of the crystallization at TI, this experiment unambiguously reveals a delay in the apparition of the secondary endotherm at TII, confirming their different physical origins. In addition, the corresponding rheological temperature sweeps presented in Figure 9b confirm the presence of several populations of crystallites by exhibiting successively two sudden falls of the storage modulus at TII and TI as already observed in homopolymers.62 This important result, supporting the isothermal analysis performed through the use of both rheology G′(t) and calorimetry X(t), reinforces the idea of a bimodal network’s formation. Nevertheless, in spite of this important progress, several questions still remain on the concomitance of the two processes as well as on a clear picture describing the interactions between adjacent crystallites. On the basis of the rich set of experimental results developed in section 3, we propose therefore to address these complex issues in the following discussion.

X(t )/X inf = 1 − exp( −t /ta)n

(7)

with ta = K−1/n being the Avrami time corresponding to 63% of the maximal crystallinity. Because n and K were found to vary with both Tiso and the TPE composition, ta was calculated for each specific case and reported in Supporting Information section 2.2. for clarity. Then, a similar procedure was applied for the neat PBTs with the raw parameters (n and K) extracted from Kulshreshtha et al.33 and Wu et al.,32 respectively, for spherulitic and lamellar growth modes. (Note that ta may have been replaced with t0.5, providing unchanged conclusions.) The results corresponding to the whole set of PBT architectures are presented in Figure 10. As expected, they emphasize a faster crystallization of the spherulitic PBT made of short chains, followed by its lamellar long chain counterpart, and at last the

4. DISCUSSION In this last section we aim at confronting all our experimental measurements with the trends generally observed in PBT based homo- and copolymers. The discussion is centered on three subtopics: (i) the influence of the PBT chain’s architecture on its crystallization, (ii) the gelation process, and (iii) the structure−properties relationship once the crystallization completed. 4.1. Overall Crystallization Kinetics of PBT in Various Environments. As evidenced all along in section 3, the chain composition and Tiso are found to play a major role in the HSs’ crystallization kinetics. As a matter of fact, the crystallization rate was systematically found to slow down exponentially with increasing Tiso, being further delayed when decreasing the PBT content within the chain from 65 to 30 wt % (see Figures 6d and 7d). Remarkably, the same kinetics were observed for HS30 and HS40. Besides, the Avrami analysis resulted in identifying a lamellar or ribbon-like growth of the crystallites (1 < n < 2) in the three TPEs. While this important characteristic is quite far from what was observed in a neat PBT of similar molecular weight (23 kg mol−1) characterized by a spherulitic growth (3 < n < 4),32 it is however rather close

Figure 10. Reconstructed overall crystallization kinetics for the PBTs in different environments (long and short neat PBTs as well as the three TPEs studied in this work). The long and short neat PBT curves were recalculated from the crystallization kinetics determined respectively in refs 32 and 31. I


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Figure 11. (a) tgel/ta as a function of ΔT for HS30 and HS65. (b) Xgel/Xinf as a function of ΔT for HS30 and HS65. Squares are extracted from tgel taken as the G′(t) − G″(t) crossover and reported in Figures 6a and 6c accordingly. Crosses are calculated for the two TPEs from eq 16 in ref 40 (Winter et al., theoretical model).

the strong impact of the chain composition on the gelation/ crystallization kinetics. Gelation in HS30. In HS30, the ratio tgel/ta is always larger than 1 and quickly drops when ΔT increases. This trend unambiguously indicates that the growth of crystallinity is predominant over the gelation, i.e., that the crystallites are rather formed locally and have time to grow before to interact significantly with each other, like e.g., through the entanglements of the soft phase. The pronounced fall of tgel/ta is understood as the raise of the nucleation kinetics leading to a higher fraction of nascent ribbons at early stages, not contributing much to the network formation. This result is corroborated by the unchanged value of Xgel/Xinf (≈1), showing that the gelation is systematically reached at later stages of the crystallization, in agreement with the low amount of HSs in the chains (30 wt %) and with the actual low crystallinity measured at long time Xinf ≈ 8 wt %.39 Based on this picture, the evolution of G′ at t > tgel (Figures 3a and 4a) is therefore not assigned to a growing crystal phase but rather to a rearrangement, or equilibration, of the so-formed network. Finally, we believe that the presence of a G″ maximum right after tgel, only visible in HS30 (see Figure 4a and Supporting Information section 1.5), is another signature of this structural reorganization at constant crystal fraction. Gelation in HS65. Despite a similar trend of tgel/ta in HS65, the corresponding values are significantly smaller than in HS30, passing even below 1 at ΔT = 32 °C (Figure 11a). Moreover, Xgel is found to strongly decrease with increasing ΔT, suggesting that the nascent crystallites have in this case the possibility to interact with each other earlier, forming a gel, and keep on growing significantly at longer time. It is worth to note that while the theoretical predictions of Winter et al.40 (based on the exponential decrease of ta and tgel with ΔT) are found to be in good agreement with our experimental values for HS30, it is not the case for HS65. In the latter sample, we believe that the very different chain architecture (PBT homopolymer versus TPE) and the thermal instabilities of the rheology experiment at short time are at the origin of this mismatch. In fact, the undershoot of the temperature (detected in Supporting Information section 1.1) is likely to accelerate (nonisothermally) the gelation of the TPE, which results in an underestimation of Xgel from Figure 6c. All the above-mentioned trends appear even clearer when the storage modulus (from rheology) is plotted against the relative crystalline ratio (from DSC), as proposed in Figure 12.

TPEs. Here, the sped up kinetics of the short PBT is assigned to the enhanced mobility of the chains allowing the crystal to grow in three dimensions, whereas the long PBT’s sluggish dynamics is believed to restrict its crystallization along two directions. Moreover, the addition of PTHF segments diluting somehow the hard units in TPEs is thus expected to further slow down their association from a statistical point of view. Beyond these basic considerations, the HS65 TPE exhibits strikingly quasi-identical crystallization kinetics as the long neat PBT. This indicates therefore, that the lengthening of the chain or its interruption by SSs results in a similar crystallization delay as well as a same way of growing, in spite of a different maximal content in crystallite. This important observation suggests also that the average length of the amorphous interphase in neat PBTs must be comparable with the SS one’s in the HS65 TPE, i.e., close to 1 kg mol−1. Following this idea, one can then apprehend the additional delay observed in both HS30 and HS40 TPEs based on their longer SSs made of 2 kg mol−1 PTHF. In fact, a SS having a larger molecular weight than the “natural” interphase in neat PBT would ineluctably perturb the chain folding, i.e., slow down the crystallization kinetics. Finally, because in this representation HS30 and HS40 exhibit indistinguishable crystallization kinetics, we believe that the length of the SS dominates the network building mechanisms over the HS concentration effect. 4.2. Impact of the Chain Composition on the Gelation Mechanism. Because we showed that the network topology was most likely different between HS65 on one side and HS30 and HS40 on the other side, we propose here to focus on the structure−properties relationship along the crystallization of the TPEs through the investigation of the gelation mechanisms. As described by Winter and coworkers,40,41 the crystallization can be characterized by two contemporaneous processes: (i) the radial growth of aggregates and fibrous strands that are responsible for the formation of the interconnected-chains network (gelation; see e.g. Figure 4a); (ii) the inner crystallization of the so-formed aggregates, i.e., their ordering along with time, mainly responsible for the raise of the crystal phase within the material. Here, it is convenient to quantify these two processes through tgel and ta. It allows then to establish the dominating process by analyzing either the variation of the ratio tgel/ta or the corresponding crystal fraction at the gelation point Xgel = X(tgel) as a function of ΔT. The corresponding results are presented in Figure 11a,b for HS30 and HS65, emphasizing J


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for HS40 and 47 °C for HS30.) For the sake of completeness, we also provide a similar comparison at ΔT = 37 °C in Supporting Information section 3, Figure S11. To summarize, while the gelation happens at the end of the crystallization in HS30, additional interactions are formed at later stages in HS40 and even more in HS65. Because it occurs at roughly 6 wt %, the gelation is undoubtedly assigned to the formation of small crystallites spanning homogeneously the whole TPE. At this point, they are connected with each other through the “usual” polymers entanglements (M̅ n ≫ Me ≈ 4 kg mol−1).19 At later stages, the crystallites have then the possibility to further grow and get closer, which makes possible their interconnection through a single (unentangled) strand, so-called “bridge” hereafter. This transition was recently predicted at ca. 15 vol. % (12 wt %) in crystallites, allowing to describe the static reinforcement in similar (T4T/PTHF) segmented copolymers.19 This major issue is discussed in detail in the last section. 4.3. Static Approach of the Structure−Properties Relationship. Related to the previous section, we finally discuss the evolution of the steady-state plateau modulus G′inf as a function of the corresponding crystallite content Xinf for the three TPEs. The result is presented in Figure 13. As one

Figure 12. Storage (squares) and loss (solid lines) moduli measured from rheology (Figure 3) plotted against the relative crystalline ratio obtained from DSC (Figure 6) for HS30 (140 °C), HS40 (160 °C), and HS65 (170 °C). Data have been shifted vertically for clarity. Arrows indicate the gel points. Rheological data in the gray zone are questionable because of thermal instabilities.

This graph notably highlights that for a given ΔT (= 44 ± 3 °C), the chain composition impacts dramatically the relative position of the gel point, i.e., the relative fraction of crystal leading to the formation of an infinite physical network. It passes indeed from Xgel/Xinf = 0.86 to 0.57 and 0.25 respectively for HS30, HS40, and HS65 (see arrows in Figure 12). However, it is striking to note that the absolute value of Xgel is found to be almost equal in the three TPEs, corresponding to 5.8, 6.7, and 5.2 wt % in crystallites (Xinf values are extracted from Figure 8d). This unexpected result evidences in consequence the unchanged character of the gelation origin, regardless of the chain composition. It suggests by extension that the HS content impacts thus mainly (i) the “temporal” crystallization kinetics, i.e., the time required for the HSs to find partners to aggregate with obviously related to their number density, and (ii) the possibility for the crystallites to access extra levels of interaction at the latest stages of the crystallization (once the material is gelified). Using the analogy of a movie to represent the whole crystallization “story” in the three TPEs would result in the following. HS30 would correspond to a movie played in slow motion, reaching the main intrigue (the gelation) after only several hours. HS40 would then correspond to the same movie played at normal speed, in which the main intrigue is discovered after a few minutes and thus lets the audience watch the full movie in a couple of hours. Lastly, HS65 would correspond to that movie played in a fast forward mode, allowing to reach the main intrigue after a few seconds and the end of the movie in a few minutes. In this last situation, the audience would still have the time to watch the hidden scenes (i.e., the enhancement of the structure at the latest stage of crystallization). In all these cases, we want to insist on the fact that the movies, i.e., the mechanisms of crystallization at early stages, are qualitatively identical. This hypothesis, suggested by the quasi-identical values of Xgel in the three TPEs, is remarkably supported by the similar storage and loss moduli at the gel point G′(Xgel/Xinf) = G″(Xgel/Xinf) in the three TPEs. They correspond to 6.4 × 104, 7.4 × 104, and 1.3 × 104 Pa respectively for HS30, HS40, and HS65, corroborating the idea of similar structures at the gel points in all our materials. (We believe that the slightly lower value reported for HS65 is due to the lower ΔT used for the experiment, 42 °C instead of 45 °C

Figure 13. Steady-state plateau modulus (G′inf) as a function of the steady-state weight fraction in crystallite (Xinf) for the three TPEs. Each point corresponds to a different ΔT. Inset: reduced reinforcement as a function of the volume fraction in crystallites (Φ) calculated as χ = G′inf/G′pthf, where G′pthf is approximated to 0.45 MPa (T > 110 °C).19 The red and blue solid lines correspond respectively to χ = exp(AΦ) and χ =

R ed *2 nHR2

Φ2 .

may have anticipated, the global trend is an enhancement of the plateau modulus with increasing the crystal fraction, revealing a strong upturn above Xinf = 15 wt %. For each TPE, the different data points correspond in reality to different ΔT extracted from Figures 2d and 7d, showing a quasi-linear trend between the rubbery plateau and the crystal fraction before an eventual saturation for the highest values of ΔT (see in particular HS40). Once again, while HS30 and HS40 exhibit a similar behavior, a much steeper slope is observed for HS65 undoubtedly related to the confinement of the SSs described in our previous work.19 More quantitatively, the inset in Figure 13 represents the “reinforcement” (defined as χ = G′inf/G′pthf) as a function of the volume fraction in crystallites (Φ). It notably allows to rationalize the material’s hardening through topological arguments as proposed in ref 19. In this article we proposed an original method to quantify the plateau modulus of K


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Macromolecules semicrystalline segmented block copolymers from the dimensions of the ribbon crystallites they contain. The main idea consists of counting the number of aggregated HSs (Nce) in a section of crystallite having a size corresponding to an “entanglement volume” in the neat soft phase (Ve).63 This number is then used to calculate the extra density in topological constraints in the hard phase following the rubber elasticity theory (νx ∼ Nce). Finally, by considering the volume fraction of the two phases (Φ and 1 − Φ) as well as their corresponding topological densities (ν and νx), we can calculate the material’s modulus and therefore the resulting reinforcement with respect to the soft matrix. Two important expressions were obtained from this work (eqs 7 and 17),19 describing respectively the TPEs modulus at low and high content in HSs.

(

Remarkably, fitting the former equation χ =

R ed *2 nHR2

)

Φ2 to

the experimental data up to (Φ = 0.12) provides an intercrystallites distance of about d* = 18 nm, in fair agreement with SAXS experiments performed on HS30 (23 nm) and HS40 (24 nm) samples (see Supporting Information section 4). The other parameters were chosen in the following way: Re = 4.4 nm, the PTHF tube diameter from ref 19; R = 0.5 nm, an approximation of the cell parameters (see b and c in α-PBT31); H ≈ 5 nm, the length of the HSs, estimated from SAXS data modeling in ref 39. In addition, the empirical simplification of the above-mentioned eq 17 (valid at high-HS content), resulting in χ = exp(AΦ), provides A = 16.97, similar to the value determined for T4T/PTHF TPEs,19 A = 18.38. This encouraging resemblance indicates therefore a possible universal behavior of the reinforcement in segmented block copolymers that one can describe from the morphology of the hard phase they contain.

Figure 14. Schematic representation of the different experiments performed in this article highlighting the two-step crystallization of the HSs. (a, b) Rheological and DSC time sweeps performed at T = Tiso. (c, d) Rheological and DSC temperature sweeps revealing the melting of different group of crystallites. Symbols I and II stand for the “primary” and “secondary” (or “hindered”) crystallizations, respectively.

fundamental evidence allowing to understand the evolution of the chains’ topology during the crystallization of polymers. For example, we believe that the high content in HSs and the resulting fast kinetics in HS65 generate a high amount of straight SS bridges at the latest stages of its crystallization. In this case, the too short SSs (1 kg mol−1) do not have the possibility to bend (nor to crystallize), resulting in the formation of highly reinforcing “crystallite−SS−crystallite” sandwiches. On the contrary, the slower association of more distant HSs leading to a delayed gelation limits the intercrystallites’ interaction to the SSs’ entanglements. In this situation, longer SSs (2 kg mol−1) preserve their ability to bend (crystallizing at low temperatures) and form entanglements with their surrounding counterparts.

5. CONCLUSION In conclusion, we have addressed jointly the crystallization and gelation kinetics of three industrially relevant PBT/PTHF segmented copolymers by using rheology and calorimetry techniques resolved in both time and temperature. As one may have anticipated, increasing the chain rigidity (i.e., the HS content) is found to increase significantly the phase separation kinetics, leading notably to an earlier gelation in HSs-rich TPEs. However, observing the same data as a function of the crystallization advancement (and not time) strongly suggests the presence of similar mechanisms in the whole set of TPEs. At the earliest stage, the HSs start to aggregate to form small crystallites that will then interact to each other through the soft-phase entanglements network, resulting in the gelation of the system. After the gelation, the crystallites keep on growing up to different levels according to the chain rigidity; the higher the HS content is, the greater the reinforcement. Finally, they will start to interact with each other via shorter polymer strands, possibly down to a single SS, limiting dramatically the mobility of residual HSs, which will thus aggregate into metastable crystallite (secondary crystallization). This scenario was evidenced through multiple experiments that we schematically represent in Figure 14 for more clarity. While in the Figures 14a and 14b I and II refer respectively to the primary (“natural”) and secondary (“hindered”) crystallizations of the HSs, they are related to the melting of the corresponding crystallites in Figures 14c and 14d. Beyond their practical interest for industrial applications, we are convinced that these observations are likely to bring new

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b01689. (1) Rheology: temperature profiles during the time sweeps, frequency sweeps performed in parallel with the isothermal crystallization, two-mode exponential fitting of the time sweeps, G″(t) data showing a maximum in HS30 and multiwave frequency sweeps evidencing flat loss factor; (2) Calorimetry: isothermal crystallization and corresponding Avrami analysis for HS30 and HS40, DSC scans after time sweeps experiments, determination of the Avrami times, Hoffman−Weeks plot; (3) Correlation between rheology and calorimetry: comparison of the characteristic times and modulus plotted versus relative crystallinity; (4) Small-angle X-ray scattering: data for the three TPEs crystallized at different Tiso in the rheometer and the corresponding fitting curves, comparison with TPEs processed through hot pressing and subsequent quenching in air as well as solvent casting (PDF) L


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(12) Huang, Y.; Chang, R.; Han, L.; Shan, G.; Bao, Y.; Pan, P. ABAType Thermoplastic Elastomers Composed of Poly(ε-caprolactoneco-δ-valerolactone) Soft Midblock and Polymorphic Poly(lactic acid) Hard End blocks. ACS Sustainable Chem. Eng. 2016, 4, 121−128. (13) Bandara, U.; Droescher, M. The two-phase structure of segmented block copoly (ether ester). Colloid Polym. Sci. 1983, 261, 26−39. (14) Zhang, C.; Zhang, Q.; Tonghui, H.; Jiang, T. Effect of Catalyst on Preparation and Properties of PBT/PTMG Poly(ether ester). Chin. J. Mater. Res. 2014, 28, 781−786. (15) Cella, R. J. In Morphology of segmented polyester thermoplastic elastomers, J. Polym. Sci.: Polym. Symp.; Wiley Online Library: 1973; pp 727−740. (16) Zhu, L. L.; Wegner, G. The morphology of semicrystalline segmented poly (ether ester) thermoplastic elastomers. Makromol. Chem. 1981, 182, 3625−3638. (17) Seymour, R.; Overton, J.; Corley, L. Morphological characterization of polyester-based elastoplastics. Macromolecules 1975, 8, 331−335. (18) Gabriëlse, W.; Soliman, M.; Dijkstra, K. Microstructure and Phase Behavior of Block Copoly(ether ester) Thermoplastic Elastomers. Macromolecules 2001, 34, 1685−1693. (19) Baeza, G. P. The Reinforcement Effect in Well-Defined Segmented Copolymers: Counting the Topological Constraints at the Mesoscopic Scale. Macromolecules 2018, 51, 1957−1966. (20) Baeza, G. P.; Sharma, A.; Louhichi, A.; Imperiali, L.; Appel, W. P. J.; Fitié, C. F. C.; Lettinga, M. P.; Van Ruymbeke, E.; Vlassopoulos, D. Multiscale Organization of Thermoplastic Elastomers With Varying Content of Hard Segments. Polymer 2016, 107, 89−101. (21) Xiong, B. J.; Lame, O.; Chenal, J. M.; Rochas, C.; Seguela, R.; Vigier, G. Amorphous Phase Modulus and Micro-Macro Scale Relationship in Polyethylene via in Situ SAXS and WAXS. Macromolecules 2015, 48, 2149−2160. (22) Makke, A.; Lame, O.; Perez, M.; Barrat, J.-L. Influence of Tie and loop molecules on the mechanical properties of lamellar block copolymers. Macromolecules 2012, 45, 8445−8452. (23) Seguela, R. Critical review of the molecular topology of semicrystalline polymers: The origin and assessment of intercrystalline tie molecules and chain entanglements. J. Polym. Sci., Part B: Polym. Phys. 2005, 43, 1729−1748. (24) Avrami, M. Kinetics of phase change. I General theory. J. Chem. Phys. 1939, 7, 1103−1112. (25) Avrami, M. Kinetics of phase change. II transformation-time relations for random distribution of nuclei. J. Chem. Phys. 1940, 8, 212−224. (26) Hillier, I. Modified avrami equation for the bulk crystallization kinetics of spherulitic polymers. J. Polym. Sci., Part A: Gen. Pap. 1965, 3, 3067−3078. (27) Velisaris, C. N.; Seferis, J. C. Crystallization kinetics of polyetheretherketone (PEEK) matrices. Polym. Eng. Sci. 1986, 26, 1574−1581. (28) Verhoyen, O.; Dupret, F.; Legras, R. Isothermal and nonisothermal crystallization kinetics of polyethylene terephthalate: Mathematical modeling and experimental measurement. Polym. Eng. Sci. 1998, 38, 1594−1610. (29) Winter, H. H.; Mours, M. Rheology of polymers near liquidsolid transitions. In Neutron Spin Echo Spectroscopy Viscoelasticity Rheology; Springer: 1997; pp 165−234. (30) Vilgis, T.; Winter, H. Mechanical selfsimilarity of polymers during chemical gelation. Colloid Polym. Sci. 1988, 266, 494−500. (31) Yokouchi, M.; Sakakibara, Y.; Chatani, Y.; Tadokoro, H.; Tanaka, T.; Yoda, K. Structures of two crystalline forms of poly (butylene terephthalate) and reversible transition between them by mechanical deformation. Macromolecules 1976, 9, 266−273. (32) Wu, D.; Zhou, C.; Fan, X.; Mao, D.; Bian, Z. Morphology, crystalline structure and isothermal crystallization kinetics of polybutylene terephthalate/montmorillonite nanocomposites. Polym. Polym. Compos. 2005, 13, 61.

AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]. ORCID

André de Almeida: 0000-0002-2731-0944 Guilhem P. Baeza: 0000-0002-5142-9670 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

All the authors are thankful to EUSMI (European Soft Matter Infrastructure - https://eusmi-h2020.eu/) for funding the rheological experiments, and to Daniele Parisi and Dimitris Vlassopoulos for their support at FORTH-IESL. The authors are indebted to the “Microstructure Technological Center” (CTμ) of University of Lyon for the access to the ultramicrotomy facility. G.P.B. thanks Nancy Eisenmenger and Carel Fitié (DSM Ahead) for providing the TPEs. A.d.A and G.P.B. express their gratitude to Ahmad Patel, who helped with the DSC experiments. A.d.A. and G.P.B. are thankful to “Programme Avenir Lyon St-Etienne” (PALSE) from the University of Lyon within the framework of “Investissement d’Avenir” (ANR-11-IDEX-0007). Partial support was provided by the E.U. FP7 (ETN Supolen GA-607937). M.N. is grateful to the French Minister of Higher Studies and Research (MESR) for his Ph.D. grant.

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