Interrelations Between Side Chain and Main Chain Packing in

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Interrelations Between Side Chain and Main Chain Packing in Different Crystal Modifications of Alkoxylated Polyesters Gaurav Kumar Gupta, Varun Danke, Tamoor Babur, and Mario Beiner J. Phys. Chem. B, Just Accepted Manuscript • DOI: 10.1021/acs.jpcb.7b00928 • Publication Date (Web): 13 Apr 2017 Downloaded from http://pubs.acs.org on April 13, 2017

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The Journal of Physical Chemistry

Interrelations Between Side Chain and Main Chain Packing in Different Crystal Modifications of Alkoxylated Polyesters Gaurav Gupta,† Varun Danke,‡ Tamoor Babur,‡ and Mario Beiner∗,‡,† Fraunhofer Institut für Mikrostruktur von Werkstoffen und Systemen IMWS, Walter-Hülse-Straße 1, D-06120 Halle (Saale), Germany, and Martin-Luther-Universität Halle-Wittenberg, Institut für Chemie, D-06099 Halle (Saale), Germany E-mail: [email protected]

∗

To whom correspondence should be addressed Fraunhofer Institut für Mikrostruktur von Werkstoffen und Systemen IMWS, Walter-Hülse-Straße 1, D-06120 Halle (Saale), Germany ‡ Martin-Luther-Universität Halle-Wittenberg, Institut für Chemie, D-06099 Halle (Saale), Germany †

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Abstract Interrelations between main chain and side chain packing are studied in a series of comb-like poly(1,4-phenylene-2,5-n-dialkyloxy terephthalate)s (PPAOTs) with C = 6 to 12 alkyl carbons per side chain by x-ray diffractometry. Two different polymorphic states, called modification A and modification B, are observed depending on thermal treatment and side chains length. For PPAOTs with short side chains (C ≤ 8) modification A is commonly observed. ’As synthesized’ (solution crystallized) PPAOTs with longer side chains (C ≥ 10) contain majorly modification B while modification A is growing during melt cooling. A solid-solid transition from modification B to modification A is observed above ≈ 70◦ C for the decyl member (C=10, PPDOT) while modification A converts under ambient conditions slowly to modification B. This indicates that modification B is thermodynamically stable at low temperature while modification A is stable at higher temperature. Crystallographic analysis shows that both modifications are characterized by an orthorhombic unit cell and a long-range ordered layered structure with alternating main chain and alkyl nanodomains. This is confirmed by 2D diffration data for shear oriented PPDOT fibers. Main difference between both modifications is the packing of the side chains which are in a crystalline state for modification B but disordered for modification A. This is concluded from values for the volume per CH2 unit VCH2 in alkyl nanodomains calculated without further assumptions from the obtained lattice parameters. Interestingly, the crystalline packing of the side chains in modification B leads to a significant increase (≈ 20%) in the application relevant π − π spacings within the main chain domains as compared to modification A. It is argued that structure formation process and thermodynamic equilibrium in comb-like polymers might be strongly influenced by a competition of the individual packing tendencies of main and side chains.

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Introduction Comb-like polymers with rigid backbones and flexible side chains are an important class of functional materials with applications in various fields like organic semi-conductors, 1,2 light emitting diodes 3 and light weight components in high performance fiber composite materials with excellent mechanical properties. 4,5 The side chains (containing long methylene sequences) are usually incorporated to enhance solubility and processability of these comb-like polymers which otherwise are difficult to attain at standard conditions. A common feature of such comb-like polymers is the formation of layered structures with typical spacings in the 1-3 nm range wherein the side chains aggregate to form alkyl nanodomains (Figure 1). This self assembling phenomenon is called nanophase separation and is found in amorphous and liquid-crystalline polymers with different side chain lengths. 6,7 The packing behavior of side chains within these domains can vary depending on the microstructure of the rigid backbones, the constraints under which structure formation occurs as well as thermal history. 7–9 The rigid backbones contain in many cases ring-like units which tend to form stacks. The resulting long range order within the main chain domains is often determining the performance of the materials. However, the packing of the methylene sequences within the alkyl nanodomains is of crucial importance for the overall packing state. Various comblike polymers have been investigated to determine the crystallinity and packing behavior of side chains. 10–16 Typically, the crystallographic analysis of comb-like polymers is made based on models considering interdigitation and tilting of alkyl side groups assuming that they are in a totally stretched all trans state (Figure 1a). 10,13,16,17 Whether or not this assumption is really applicable seems to be open and is controversially debated in many cases. Ballauff et al. investigated a series of alkoxylated polyesters and found that depending on the side chain lengths and thermal history (solvent precipitation) PPAOTs can exist in three different long range ordered polymorphs namely, the modifications A, B and C. 18 While modification C is only found in samples with C ≤ 4 alkyl carbons per side chain, modifications A and B are most relevant for samples with longer alkyl groups. A major difference between modifications 3



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very detailed models or was made based on structure simulations depending on additional assumptions. 24 In this context achieving orientation in PPAOTs can serve two purposes in parallel - it is not only crucial for obtaining materials with optimized mechanical properties but it is also important for reaching a better understanding of crystal structure. In this work, the structure of comb-like poly(1,4-phenylene-2,5-n-dialkyloxy terephthalate)s with rigid backbone and different side chain length (C = 6 to 12) is analyzed in detail. Temperature-dependent x-ray diffraction measurements on isotropic samples after different thermal treatment are combined with 2D scattering measurements on oriented fibers obtained by ram extrusion. Main focus of this work is to study interrelations between packing of side chains and main chains in different layered polymorphic states of PPAOTs, namely modifications A and B.

Results Temperature-dependent x-ray scattering profiles for PPOOT (C = 8) and PPDOT (C = 10) measured during stepwise first heating and cooling scans at temperatures between T = 30 ◦

C and T = 230 ◦ C are provided in Figure 2 as representative examples. The temperature-

dependent scattering profiles for PPHOT (C = 6) and PPDDOT (C = 12) with different alkyl side chain lengths show qualitatively similar behavior (see Supplementary information, Figure A1). A main feature in the x-ray diffraction patterns for PPOOT (Figure 2 a) is a strong reflection at a scattering vector q of 3.87 nm−1 (labeled as qA 100 ) indicating a long range ordered state at the lowest measurement temperature (T = 30 ◦ C). A lamellar structure on the nanoscale is indicated by the presence of higher order reflections occurring A A at qA 200 and q300 . The relevant inter-planar spacing is d100 ≈ 16.2 Å as calculated from the

peak position of the first order peak qA 100 using Bragg s equation, d = 2π/q. This lamellar ′

structure corresponds to a periodic arrangement of nanodomains formed by rigid backbones (comprising of aromatic rings) and aggregated side chains. The obtained spacing is in good

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qA100

qA100

C=8

(a)

qA020

qAxxx qA 200

o

30 C

qA300

qA020

o

30 C

Intensity (a.u.)

Intensity (a.u.)

qAxxx qA 200

C=8

(b)

o

230 C 2

4

6

8

10 12 14

o

16 18 20

2

4

6

(c)

8

10 12 14

16 18 20

230 C

scattering vector q (nm-1)


qA100

(d)

C=10

C=10

qA100 qB100 qAxxx qA200 qA 300

o

30 C

qA020 30 oC

Intensity (a.u.)

Intensity (a.u.)

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o

o

230 C 2

4

6

8

10 12 14

230 C

16 18 20

2


4

6

8

10 12 14

16 18 20


Figure 2: Temperature-dependent scattering profiles of PPOOT (a,b) and PPDOT (c,d) during the first heating (a,c) and first cooling (b,d) scan. The scan rate was 10 K/min. The Miller indices of the important reflections are indicated. The heating scan of PPDOT (c) shows a solid-solid phase transition in the temperature range T = 70 - 120 ◦ C from modification B to modification A which is irreversible during cooling. Such a phase transition is absent for the PPOOT sample.

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agreement with previously reported values for modification A in PPOOT. 17 The reflection −1 indexed as qA corresponds to an interplanar spacing dA 020 at q ≈ 17.2 nm 020 of about 3.65 Å

and is related to the distance between the aromatic rings within the stacks of main chains −1 (π − π spacing). In addition, a peak labeled as qA is clearly seen in the xxx at 5.3 nm

diffraction data of modification A for all investigated PPAOTs (Figures 2 and A1). 19 Results for oriented samples presented below confirm that the related spacing dA xxx of about 11.9 Å corresponds to the periodicity (or side chain-to-side chain distance) in main chain direction. Note that this distance between neighbored side chains was estimated earlier to be about 12.6 Å. 23,24 During heating, the Bragg peak caused by (020) lattice planes seems to vanish at a lower temperature than the reflections related to (100) lattice planes. This implies that a significant part of the polymer is still long range ordered although the stacking of backbones is no longer observed. At high temperatures T ≥ 220 ◦ C a conventional amorphous halo in the wide angle scattering range 10 nm−1 < q < 18 nm−1 and a broad ’pre-peak’ with a peak position close to the (100) peak at q ≈ 3.5 nm−1 occurs for PPOOT indicating that the nanophase separation of main and side chains is also preserved to a certain extent in the molten state without long range order. 6 In contrast to PPOOT, the scattering patterns measured during heating on as synthesized PPDOT (Figure 2 (c)) show interestingly some other features. A strong reflection at q = ◦ B 4.6 nm−1 (labeled as qB 100 ) is observed at T = 30 C together with higher orders at q200 −1 and qB < q < 16.5 nm−1 ) a relatively broad peak 300 . Further, in the WAXS range (14 nm

being a superposition of two peaks is observed. A comparison of the position of the (100) reflection with literature data indicates clearly that the as synthesized PPDOT sample occurs as modification B. Both, the qB 100 peak as well as the broad peak at larger q become weaker with increasing temperature and vanish completely in the temperature interval 70 ◦ C < T < 120 ◦ C. This is illustrated based on selected scattering profiles in Figure 3. Simultaneously, in the same temperature range, a Bragg reflection at q ≈ 3.3 nm−1 (labeled as qA 100 ) develops together with a peak at q ≈ 17.3 nm−1 (marked by a dotted red line in Figure 2 (c)). These

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qA100 qB100 qA 200

qA020

qB200

qA100

qA020

qAxxx qA200 Intensity (a.u.)

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qAxxx qA

qA020

200

RT 3 years storage T = 30 oC melt cooled T = 140 oC

qA100 qB 100 qA200 qB 200

T = 90 oC

qB100 qB200 2

T = 30 oC

4 6 8 10 12 14 16 18 20 scattering vector q (nm-1)

Figure 3: Scattering profiles of PPDOT at selective temperatures during the first heating −1 run exhibiting transition from modification B to A. While the peak qB 100 (q = 4.6 nm ) is B observed at T = 30 ◦ C, qA 100 evolves at higher temperatures. No q100 is observed at T = 140 ◦ C where modification A is the stable state. Changes in the WAXS region are also observed where the broad peak(s) (14 nm−1 < q < 16.5 nm−1 ) at T = 30 ◦ C gradually disappear −1 and qA ) is prominent at T = 140 ◦ C. The scattering profile for a PPDOT 020 (q ≈ 17 nm sample cooled from the molten state to 30 ◦ C containing only modification A is given for comparison. The top most scattering profile, obtained after prolonged storage under ambient conditions of a melt cooled PPDOT sample, exhibits the partial conversion of modification A to B.

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reflections belonging to modification A of PPDOT dominate the diffraction patterns at higher temperatures until the reflection qA 100 finally disappears when the polymer melts completely A at T = 230 ◦ C. The inter-planar spacing corresponding to qB 100 and q100 are calculated to

be 13.68 Å and 18.98 Å, respectively. These values are in good agreement with spacings reported for modification B and modification A of PPDOT in the literature. 17 An estimate of the ratio of fraction A / fraction B based on the estimated peak areas under the (100) peak reveals an increase from 0.58 at T = 80 ◦ C to 2.9 at T = 100 ◦ C. This confirms that within this temperature range, 70 ◦ C < T < 120 ◦ C, the PPDOT sample is undergoing a solid-solid transition from modification B to modification A. This is also supported by a strong endothermal peak occurring in DSC heating scans in this temperature range for an as synthesized PPDOT sample containing only modification B which is missing in case of samples containing only modification A (cf. SI, Figure A2). Note that a similar transition was also observed in the scattering pattern for PPDDOT samples during the first heating scan (cf. SI, Figure A1 (c)). Interestingly, the first cooling scans of all the investigated samples do not show any reflection corresponding to modification B (Figure 2 (b, d) and SI Figure A1(b,d)). During cooling modification A is commonly growing below a certain temperature near 200 ◦ C which is in detail system-dependent. The scattering pattern measured at T = 30 ◦ C after cooling from the melt state are compared for all the members of PPAOTs in Figure 4 (a). The peak at q ≈ 3.3 nm−1 corresponds to the (100) reflection from the lattice planes of modification A. The other reflections corresponding to the (xxx) and (020) lattice planes are also observed and labeled. The reflection from the (100) lattice planes shifts systematically to lower q values with increasing side chain lengths indicating an increased interplanar spacing d100 . The increase in dA 100 with increasing side chain lengths (Figure 4 (b)) implies an increased thickness of the alkyl nanodomains as commonly reported for other comb-like polymers exhibiting nanophase separation. 25,26 The interplanar spacing dB 100 corresponding to modification B for PPDOT and PPDDOT are compared in Figure 4 (b) with the corresponding spacings dA 100 for modification A. The

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obtained lamellar spacings d100 for both modifications are comparable to the experimental values determined by Ballauff et al. 17 (a)

qA100 qAxxx

qA200

Intensity (a.u.)

qA200

C=6 C=8

C = 10 C = 12

2

4

6

8

10

12

14

16

18

20

scattering vector q (nm-1) (b) 20

dA100 dA100 Ref[5] dB100

15

dB100 Ref[5]

‡

d100 ( )

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10

5

0

0

2

4

6

8

10

12

C atoms per side chain

Figure 4: (a) Room temperature scattering profiles (after cooling from melt) for a series of PPAOT samples (C = 6 to C = 12) containing a high fraction of modification A. The Miller indices for the most important scattering peaks are labelled. The vertical dotted lines at q ≈ 5.3 nm−1 and 17.3 nm−1 indicate the weak dependence of (xxx) and (020) peaks on side chain lengths. (b) Dependence of inter lamellar spacing d100 on number of carbon atoms in the side chain for modification A (filled orange circles) and modification B (filled green circles) and its comparison with values from Ref[ 5 ](open circles and squares). Based on the above observations it is clear that modification B in PPAOTs with longer side chain lengths (C ≥ 10) (formed immediately after synthesis) is converted gradually into modification A by an irreversible process similar to a solid-solid transition upon heating. Modification A is subsequently growing during cooling at high temperatures (200 ◦ C < T < 160 ◦ C) but does not transforms to modification B at lower temperatures. However, scatter10


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ing measurements on samples stored under ambient conditions for long times (Figure 3) do reveal the presence of modification B. A slow transition of modification A to modification B is indicated. Whether or not and on which time scale this is leading to a complete conversion to modification B is an aspect which needs to be investigated further. However, the fact that modification A converts isothermally to modification B of PPDOT shows that this modification is thermodynamically more stable under ambient conditions. The finding that modification A is observed during cooling of molten PPDOT down to low temperatures has most likely kinetic reasons. In order to ascertain the packing behavior of the side chains, an important aspect is to confirm that the reflection labeled as qA xxx represents the periodicity between the neighbored alkyl groups along the polymer backbone. This is already indicated by the similarity of the corresponding spacings dA xxx for all investigated PPAOT samples in Figure 4 (b) but additional measurements on oriented samples are required to support this interpretation. Figure 5 (a) shows the 2D scattering pattern and the integrated scattering profile (in the intermediate scattering range) of PPDOT fibers prepared using a home-made extruder at T = 120 ◦ C with a shear rate of 600 sec−1 as a representative example. The initial observation from Figure 5 (a) is that the lamellar morphology (modification A) present prior to the processing is preserved with almost constant inter-lamellar spacing. Furthermore, an anisotropic intensity distribution is observed with intensity maxima along the equatorial position for the qA 100 peak corresponding to the layered morphology and along the meridional position for the qA xxx peak implying that the surface normals of the respective lattice planes are orthogonal to each other. This provides a clear evidence that the peak at qA xxx corresponds to the side-chain distance along the backbone. Note that the shear induced orientation was similar for PPDOT irrespective of the chosen processing conditions. An −1 additional weak peak at qB corresponding to the inter-lamellar spacing for 100 ≈ 4.7 nm

modification B is also observed with an orientation distribution similar to that observed for modification A. Obviously a certain fraction of modification B exists in extruded PPDOT

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samples. Possibly, partial conversion of modification A to B appeared during the storage of extruded samples under ambient conditions. Further, the 2D WAXS pattern (Figure 5 A (b)) exhibits an intensity maximum along the equatorial position for the reflection at q020

≈ 17.3 nm−1 corresponding to the π − π stacking of the backbones in case of modification A. This reflection is accompanied by two other distinct reflections at q ≈ 14.3 nm−1 and q ≈ 15.3 nm−1 in equatorial position. These reflections are also observed in PPDOT samples exhibiting modification B and gradually disappear when the transition from modification B to modification A occurs (Figure 3). Interestingly, the orientation of modification A and B fractions in extruded PPDOT samples is practically identical as it would be expected in case of crystals forming by a solid-solid transition or growing from nuclei which are pre-oriented similarly in the shear field. Altogether the results from scattering measurements on extruded PPDOT fibers in Figure 5 reveal that the fibers predominantly contain modification A crystals where the polymeric backbones are oriented along the shear direction. The fact that the (100) and (xxx) lattice planes are oriented perpendicular to each other supports the argument that the (xxx) reflection is indeed related to the periodicity along the main chains like indicated in Figure 1. Henceforth, the indices (xxx) can be replaced by (001).

Discussion Crystallographic states and side chain packing An interesting finding in case of PPAOTs is that at least two different polymorphs, modification A and modification B, can appear depending on side chain length and thermal treatment. Both modifications are characterized by a lamellar morphology with alternating main chain and alkyl nanodomains. Modification A has been observed for all investigated PPAOTs while modification B has been found only in the higher members with C ≥ 10. Similar behavior with two or more polymorphs has been reported for other series of comb12


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(a)

—n ÚÙÙ

qA100 Intensity (a.u.)

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—m ÚÙÙ

—m ÙÙÚ

qB100 qA001

2

qA200

3 4 5 6 7 8 scattering vector q (nm-1)

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(b)

qA100 Intensity (a.u.)

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—m ÙÛÙ

qB100 B 200

q qA001 A q200 qA300 2

4

6

qB020 qB120 qA020

8 10 12 14 16 18 20


Figure 5: 2D image of the scattered intensity distribution of the extruded fiber processed at T = 120 ◦ C and shear rate 600 s−1 in the (a) intermediate and (b) WAXS range along with the corresponding integrated scattering profiles. The important peaks are marked. The fiber is placed in such a way that it is vertical and lying in the plane of the paper. Predominant scattering is seen at the equatorial positions for (100) peaks in modification A and B while meridional scattering is stronger in case of the (xxx) peak. The black regions in part (b) are due to blind areas of the detector.

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like polymers and seems to be a relatively general feature in case of comb-like polymers. 17 The question arises what are the crystallographic differences between different modifications and to what extent a competition of main and side chain packing is driving the occurrence of different polymorphs. 26 The crystallographic structure of both modifications has been analyzed in detail based on the x-ray diffraction pattern taken from powder-like and oriented samples. Lattice models for both modifications A and B have been derived. Figure 6 gives as an example the scattering pattern of the PPDOT sample measured at room temperature (a) after cooling from the melt (modification A) and (b) on an as synthesized specimen (modification B). The scattering pattern for both modifications are indexed assuming an orthorhombic unit cell (inset Figure 6(a)). The lattice parameters are a= 1.89 nm, b = 0.726 nm and c = 1.18 nm for modification A and a = 1.36 nm, b = 0.88 nm and c = 1.18 nm for modification B. The crystallographic densities as calculated from the volume of the unit cells and the monomeric mass of 558 g/mol are 1.14 g/cm3 and 1.3 g/cm3 considering two monomers per unit cell for modifications A and B, respectively. Note that the relative orientation of the peaks in case of shear oriented PPDOT fibers are in agreement with these lattice models (Figure 5). Similarly, peak fitting based on an orthorhombic lattice model was also possible for the other members of the PPAOTs under investigation (C = 6, 8 and 12; cf. SI Figure A3). The corresponding lattice parameters for modification A (and where available modification B) of these samples are summarized in Table 1. Despite of the commonly observed orthorhombic unit cell there are significant differences comparing the a parameters of modifications A and B for the higher members (C ≥ 10). This can be understood as a first indication for differences regarding the packing state of the side chains since the backbones are identical. A more detailed model-free analysis of the packing state of the methylene sequences in the alkyl nanodomains can be done by comparing the average volume per CH2 unit VCH2 . As recently discussed the VCH2 values can be calculated based on the lattice parameters given in Table 1 using 15

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(a)

C = 10 (Mod A) A 100

q

Y

Y Y

Intensity (a.u.)

O C O

b

Y

Y

Y

O C O

A

qA001 q200

4

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O C

alkyl nano-domain

a

2

Y O O C

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Y O O C

O C

c

qA020

A 300

q

10

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qA510

18

20

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(b)

C= 10 (Mod B)

Intensity (a.u.)

qB100

qB020 qB120 qB001 qB010 qB200 qB210

2

(c)

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14

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amorphous PE PPAOTs (mod A)

/

28

24

crystalline PE

22

‡3

16


V CH2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


rreg P3ATs AOPPVs

20 18

PPAOTs (mod B) 16 0

2

4

6

8

10

12

Alkyl carbons per side chain

Figure 6: Scattering pattern of (a) melt pressed and (b) as synthesized PPDOT samples measured at room temperature exhibiting modification A and modification B respectively along with peak indexing using an orthorhombic unit cell. Schematic of the orthorhombic unit cell with the light organge box denoting the crystalline/liquid-like alkyl nanodomain is shown in the inset. (c) Average volume per methylene unit VCH2 dependent on side chain length for rreg P3ATs (black squares), 7 PPAOTs (modification A) (red circles), AOPPVs (blue diamonds), 15 and PPDOT (modfication B) (green circle). The VCH2 values for amorphous polyethylene (dotted line) 27 and crystalline polyethylene (solid line) 28 are given for 15 comparison. All values are measured under Plus ambient conditions. ACS Paragon Environment


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Table 1: Orthorhombic unit cell parameters and crystallographic density of PPAOTs.

Modification A PPHOT PPOOT PPDOT PPDDOT Modification B PPDOT PPDDOT

C

a b c M◦ ρa (nm) (nm) (nm) (g/mol) (gm/cm3 )

6 8 10 12

1.39 1.62 1.89 2.10

0.71 0.73 0.73 0.73

1.16 1.16 1.18 1.19

446 502 558 614

1.29 1.21 1.14 1.12

10 12

1.36 1.54

0.88 0.88

1.18 1.19

558 614

1.30 1.26

a

The crystallographic density is calculated as ρ = nM◦ /Vuc NA ; where n is the number of monomers per unit cell (here n = 2), Vuc is the volume of unit cell, NA is Avogadro constant and M◦ is the monomer mass.

VCH2 =

(d100 − dmc ) × 2d020 × d001 . (4 × C)

(1)

This requires only information about the distance between the points where the side chains are attached to the rigid backbone, i.e. d001 and d020 , as well as the thickness of alkyl nanodomains dalkyl = (d100 - dmc ) with dmc being the average thickness of the main chain layers (Figure 1). These dmc values are estimated for both PPAOT modifications from a linear extrapolation of the d100 values observed for members with different side chain lengths (Figure 4 (b)). We obtained 6.7 Å for PPAOTs. Note that the linear dependence on side chain length is in this model a direct consequence of an identical volume per CH2 unit and that the methyl end groups (CH3 ) are treated in Eq.(1) like methylene units (CH2 ). The results for PPAOTs samples containing modification A reveal that the side chains are non-crystalline and disordered within the alkyl nanodomains since the VCH2 values are close to those for amorphous polyethylene (Figure 6(c)). 15 The situation in PPDOT and PPDDOT samples containing modification B seems to be quite different. In this case the volume per CH2 unit is much smaller and even below the value expected for crystalline polyethylene (Figure 6(c)). Such a high packing density of the methylene sequences implies

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a high degree of crystallinity within the alkyl nanodomains. Note that the differences between the VCH2 values for both modifications are quite large while the scatter depending on side chain length C is rather limited for a given modification. This hints to an uniform side chain packing state for a given modification. We understand the modification-dependent differences in the VCH2 value as a strong argument for the idea that crystallization of side chains drives the formation of different PPAOT polymorphs with lamellar morphology. It is obviously the packing state of the side chains that causes the big differences in the d100 values between modifications A and B in PPAOTs. Note that packing models for modifications A and B reported previously are mostly based on the assumption that the side chains are stretched (all-trans), interdigitated and mainly oriented normal to the main chains in case of modification A, while they are assumed to be tilted (with an angle of ≈ 50◦ ) in case of modification B. 5,17,24 The packing model presented here excludes the presumptions of having stretched and differently tilted side chains. The results of our modeling confirm nevertheless the idea that modification A is a ’mesophase’ 17 incorporating disordered side chains in a long range ordered lamellar structure.

Interrelations between main and side chain packing Starting from the conclusions drawn so far it seems to be interesting to consider interrelations between side chain and main chain packing. Inspecting the scattering pattern of PPDOT in Figure 6 carefully one can conclude that the (020) reflections of modification A A corresponding to the π − π stacking at q020 ≈ 17.3 nm−1 is not observed in the scattering

pattern of modification B. In modification B, the (020) reflection is shifted to lower q values B and appears at q020 ≈ 15.3 nm−1 . This evidences that the crystalline packing of side chains

in modification B causes the backbones to pack differently leading to an ≈ 20% increment in the lattice parameter b. In terms of functional properties this is a large difference in the π −π spacing causing usually significant differences in application relevant properties. This is known especially in terms of electronic properties like conductivity since larger stacking dis17



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tances substantially reduce the overlap of electronic orbitals. 14,29,30 Although not studied in detail yet, the mechanical properties should behave similarly since the interaction strengths is affected. In summary, one can conclude that the reported packing models suggest that the functional properties of comb-like polymers are to a large extent defined by the packing state of the side chain since these sub-units strongly influence the packing of the rigid backbones. This statement is obviously true although the functional parameters are majorly determined by the packing state of rigid backbones. Central message is that changes in side chain packing in comb-like polymers will always influence main chain packing and therefore (at least potentially) the performance of functional materials. It is important to consider this in cases where different polymorphs may occur although they are not so well distinguished in the discussion. An example of this type might be the famous case of regio-regular poly(3-hexyl thiophenes) [rreg P3HT] where structural states with crystalline side chains seem to exist 16 beside of states where the side chains remain disordered within alkyl nanodomains 26,31 depending on microstructure and treatment of the samples. The influence of side chain packing on the overall performance has to be also kept in mind if side chains with variable microstructure are attached to identical backbones. This a well known playground for performance optimization. Avoiding crystallization of side chains by introducing "defects" (e.g. ethyl-hexyl instead of n-octyl side chains) like done in many organic semi-conductors with high performance and comb-like architecture 32,33 can lead to robust functional states but means at the same time that improving the performance by choosing other polymorphic states is no longer possible. Otherwise, the absence of different polymorphic states can avoid problems towards long term stability which go hand in hand with the occurrence of solid-solid transitions in comb-like polymers. 7 Finally, we will shortly discuss how the interrelation between main and side chain packing considered above can be understood based on a competition of the packing tendencies of main and side chains during the crystal formation. Native packing and thermodynamic equi-

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librium state of main chains and alkyl groups in comb-like polymers are basically different. This fact seems to be of major importance for the formation of different crystalline modifications. It is well known that rigid backbones containing ring-like units commonly tend to form π − π stacks already at very high temperatures (≥150 ◦ C) 34 while free methylene sequences with limited length (C ≤ 12) form only at low temperatures crystals as equilibrium state. 35 Otherwise, it is also clear that the melting enthalpies per gram methylene unit are significantly larger than those for systems forming π − π stacks. This can be concluded from calorimetric data for various polymers reported in the literature (cf. Ref. 36 ). The structure formation behavior of comb-like PPAOTs during cooling may be a result of these properties of the sub-units. At high temperatures the rings in the backbones may form in a first step π − π stacks while the side chain remain disordered since there is no intrisic driving force in the alkyl nanodomains to crystallize. Despite of that a long range ordered lamellar morphology can exist due to nanophase-separation of main and side chain domains. This state is commonly called modification A or ’mesophase’ by Ballauff et al. 17 We assume based on the results of this study that this is at least at high temperatures (≥100 ◦ C) the thermodynamically stable state in all investigated PPAOTs. At lower temperatures, in particular close to ambient, modification B with crystalline methylene units in the side chains seems to be thermodynamically stable for higher PPAOTs (C ≥ 10). This can be concluded from (i) the appearance of modification B in solution crystallized PPDOT and PPDDOT samples (evaporated at room temperature), (ii) a transition from modification B to modification A during heating at temperatures between 70 ◦ C and 120 ◦ C and (iii) a slow growth of modification B at room temperature in PPDOT samples which contain initially only modification A. During cooling, the once formed main chain stacks of modification A do probably hinder the crystallization of methylene sequences at lower temperatures preventing the transition to modification B. Hence, modification A occurs even at low temperatures where modification B with crystalline side chains should be thermodynamically stable. Over time, the free energy difference between the amorphous and crystalline state of the methylene sequences acts

19



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as a driving force and the system tries to convert to its equilibrium state being modification B under ambient conditions. Whether or not the side chains are really tilted with respect to the main chains and to what extent the side chains are crystallized in modification B is still a question for further consideration. Very small VCH2 values point to a high degree of crystallinity. The parallel increase in d020 might be due to an optimization of the overall free energy. Alternative experimental methods like NMR 37 or FTIR spectroscopy 12 may help in obtaining deeper insights into the packing state of the methylene sequences. Despite of that, the findings of this work highlight clearly the importance of interrelations between side chain packing and main chain packing in higher members of the PPAOT series under investigation. Further consideration of this aspect should help to understand structure formation processes and polymorphic states occurring in different types of comb-like polymers.

Conclusion In the investigated series of poly(1,4-phenylene-2,5-n-dialkyloxy terephthalate)s (PPAOTs) with C = 6 to 12 alkyl carbons per side chain two different polymorphic states, called modification A and modification B, are observed depending on thermal treatment and side chains length. Cooling PPAOTs samples from the melt state leads commonly to the formation of modification A while as synthesized samples (crystallized in the presence of solvent) contain mainly modification A in case of short side chains (C ≤ 8) but mainly modification B for longer side chains (C ≥ 10). For the latter a solid-solid transition from modification B to modification A is observed during heating while prolonged storage under ambient conditions leads to a partial conversion of modification A to modification B. A closer inspection of the crystallographic states shows that both modifications A and B are characterized by an orthorhombic unit cell and a long-range ordered layered structure where main chain domains and alkyl nanodomains alternate. This is also confirmed by 2D diffraction data for a shear oriented PPDOT sample produced by ram extrusion. Major differences between both

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modifications are the unit cell dimensions and the packing state of the methylene sequences within the alkyl nanodomains. The corresponding volumes per CH2 unit VCH2 suggest crystalline packing of the side chains in modification B but disordered methylene sequences in case of modification A. This difference in side chain packing results obviously in a significant variation (≈ 20%) of the π − π spacing within the main chain layers determining functional properties. This clearly indicates interrelations between main and side chain packing. It is assumed here that the competition of the individual packing tendencies of main and side chains is very important for an understanding of structure formation processes and thermodynamic stability of different modifications in comb-like polymers.

Experimental Materials Poly(1,4-phenylene-2,5-n-dialkyloxy terephthalates)(PPAOTS) with C = (6 to 12) alkyl carbons per side chain are studied. The molecular weights and PDIs as measured by GPC using polystyrene standards are given in Table 2. The melting temperatures of the polymer has been estimated from temperature dependent x-ray measurements. Details about the synthesis of these alkoxylated polyesters are described elsewhere. 19 Table 2: Structural parameters of PPAOTs C PPHOT 6 PPOOT 8 PPDOT 10 PPDDOT 12

Mw (kg/mol) 183 78 95 5.2

PDI 1.02 6.30 6.50 2.01

Monomer mass M◦ (g/mol) 446 502 558 614

Tm (◦ C) ≈ 230 ≈ 220 ≈ 230 ≈ 170

Oriented fibers for the PPDOT member are produced using a home built ram extruder allowing processing of small amount of samples (≈ 200 mg) with defined temperature and processing speed. The polymer was fed to a cylindrical reservoir which was heated to the 21



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required processing temperature. The upper plate of the extruder unit (to which the ram is attached) was coupled to the clamps of Universal Testing Machine (Zwick) operating in compression mode with a defined processing speed. The corresponding shear rate was then calculated based on the capillary dimensions and the processing speed. Fibers using the ram extruder were prepared at three different temperatures T = 120 ◦ C, 140 ◦ C and 160 ◦ C for three different shear rates 20 s−1 , 300 s−1 and 600 s−1 .

X-ray scattering Temperature-dependent x-ray diffraction measurements in reflection mode were performed using an Empyrean diffractometer (PANalytical) equipped with the temperature chamber TTK 450 (Anton Paar). The heating and cooling rate was kept at 10 K/min. The emitted CuKα radiation is parallelized and monochromatized using a parallel beam mirror (λ = 1.54 Å). The scattered intensity passes a parallel plate collimator (0.27◦ ) and is detected by a Pixel 3D detector with 19 channels of 0.055 µm size combined to be used as a receiving slit. The scan range was 1.5 nm−1 < q < 20.0 nm−1 , the step size 0.05◦ and the counting time per step 1 s. Room temperature x-ray diffraction measurements containing mainly mesophase A were performed on melt pressed samples. The thermal treatment is described elsewhere. 19 X-ray diffraction measurement on oriented fibers are performed (i) in the intermediate scattering vector range using a pinhole instrument designed by JJ X-rays with a Rigaku rotating anode as radiation source, an Osmic multilayer optics and a Bruker Hi-star 2D detector and (ii) in the WAXS range using a micro-focus radiation source and a PILATUS detector in transmission. The samples were measured in such a way that the X-ray beam was perpendicular to the extrusion direction. The scattering vector range was 1.5 nm−1 < q < 9 nm−1 . CuKα radiation was used (λ = 1.54 Å) and the calibration was performed using a silver behenate standard.

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Acknowledgement Financial support by the DFG in the framework of the SFB/TRR102 ’Polymers under multiple constrains: restricted and controlled molecular order and mobility’ (projects A15 and B14) is acknowledged.

Supporting Information Available Temperature dependent scattering profiles of PPHOT and PPDDOT during the stepwise heating and cooling along with the unit cell fits for modification A and B for PPHOT, PPOOT and PPDDOT can be found in the supporting information. This material is available free of charge via the Internet at http://pubs.acs.org/.

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Table of Content Graphics Y

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