Mesomorphism and shape-memory behavior of main-chain liquid-crystalline co-elastomers: Modulation by the chemical composition

Article abstract of DOI:10.1021/ma501164u

The mesomorphic behavior, shape shifting, and shape-memory types of two homologous series of main-chain liquid-crystalline co-elastomers (MC-LCEs) with composition-tunable transition temperatures and mesophases structures are reported. The weakly cross-linked macromolecules were prepared in a one-step procedure by the versatile platinum-catalyzed hydrosilylation reaction between siloxane oligomers (linear spacers and netlike point molecules) and liquid-crystalline binary mixtures, composed of defined proportions of divinylic monomeric mesogens. First, the two sets of alternated, linear copolymers with disiloxane or trisiloxane chain extender, respectively, were investigated. They either exhibit both smectic C (SmC) and nematic (N) phases or a single SmC phase, the nature of which being dictated by the length of the inserted siloxane moieties, and the temperatures ranges regulated by the composition of the mesogenic mixture (mesogens ratio). As anticipated, the thermal behavior of the corresponding co-elastomers is not greatly affected by the weak reticulation (less than 5% w/w), whose transition temperatures (from the ambient to ca. 100 °C) and mesophases types (SmC and N) are quasi-replicated from the polymers and still majority governed by the intermingling of the constitutive components (mesogens ratios, siloxane, and comonomers). Moreover, these MC-LCEs exhibit two types of shape-memory behaviors, which are imposed selectively by the nature of the siloxane extender and which are understood in relation to the mesophases' structures. The combination of small-angle X-ray scattering and thermoelastic experiments gives some fundamental insights into the relationships between molecular structures and macroscopic elastic properties in MC-LCEs in general, which should be beneficial for future designs of soft materials with desired shape-memory properties.

Full text of DOI:10.1021/ma501164u

Article
pubs.acs.org/Macromolecules
Mesomorphism and Shape-Memory Behavior of Main-Chain Liquid-
Crystalline Co-Elastomers: Modulation by the Chemical Composition
Alfonso Ramon García-Marquez,† Benoît Heinrich,† Nicolas Beyer,† Daniel Guillon,†
́
and Bertrand Donnio*,†,‡
†Institut de Physique et de Chimie des Mater
Loess BP 43, 67034 Strasbourg, Cedex 2, France
́ ́
iaux de Strasbourg (IPCMS), UMR 7504 (CNRS-Universite de Strasbourg), 23 Rue du
‡Complex Assemblies of Soft Matter Laboratory (COMPASS), UMI 3254 (CNRS-Solvay-University of Pennsylvania), CRTB, 350
George Patterson Boulevard, Bristol, Pennsylvania 19007, United States
S
* Supporting Information
ABSTRACT: The mesomorphic behavior, shape shifting, and
shape-memory types of two homologous series of main-chain
liquid-crystalline co-elastomers (MC-LCEs) with composition-
tunable transition temperatures and mesophases structures are
reported. The weakly cross-linked macromolecules were
prepared in a one-step procedure by the versatile platinum-
catalyzed hydrosilylation reaction between siloxane oligomers
(linear spacers and netlike point molecules) and liquid-
crystalline binary mixtures, composed of defined proportions
of divinylic monomeric mesogens. First, the two sets of
alternated, linear copolymers with disiloxane or trisiloxane chain extender, respectively, were investigated. They either exhibit
both smectic C (SmC) and nematic (N) phases or a single SmC phase, the nature of which being dictated by the length of the
inserted siloxane moieties, and the temperatures ranges regulated by the composition of the mesogenic mixture (mesogens ratio).
As anticipated, the thermal behavior of the corresponding co-elastomers is not greatly affected by the weak reticulation (less than
5% w/w), whose transition temperatures (from the ambient to ca. 100 °C) and mesophases types (SmC and N) are quasi-
replicated from the polymers and still majority governed by the intermingling of the constitutive components (mesogens ratios,
siloxane, and comonomers). Moreover, these MC-LCEs exhibit two types of shape-memory behaviors, which are imposed
selectively by the nature of the siloxane extender and which are understood in relation to the mesophases’ structures. The
combination of small-angle X-ray scattering and thermoelastic experiments gives some fundamental insights into the relationships
between molecular structures and macroscopic elastic properties in MC-LCEs in general, which should be beneficial for future
designs of soft materials with desired shape-memory properties.
Actuation of LCEs essentially relies on the elastic response of
the material subjected to an excitation triggered by temper-
ature,3,4 chemical,3,4 light,13 or electric/magnetic fields.14 In
thermally activated systems, the disordering of the low-
dimensional liquid-crystalline phase, resulting from the
increasing collective molecular movements (mesogens), may
lend to large macroscopic deformations of the polymeric
network, and the maximum of the amplitude changes usually
occurs at the mesophase to isotropic state transformation
temperature. Prior to the observation of this remarkable
behavior and in order to ensure large actuation amplitudes
and stimuli reversibility, the LCE materials need to be
macroscopically oriented in the mesophase and must remain
stable in this conformation. Large aligned “domains” can be
obtained via two-step cross-linking reaction procedures,
involving first an intermediate orientation process of the
INTRODUCTION
■
The concept of two-component materials combining at once
entropy-induced rubber elasticity and liquid-crystalline ordering
was proposed nearly 40 years ago by de Gennes1 but was not
experimentally reached until the successful synthesis in 1981 of
the first liquid-crystalline elastomer (LCE) by Finkelmann et
al.2 The exceptional elastic behavior of LCE materials3,4 and
particularly their ability to respond to various external stimuli
by significant macroscopic deformations (shape shifting), and
to recover their original shape under appropriate excitation
(shape memory effect),5 indeed make them potentially
attractive for the implementation into smart materials,
including artificial actuators and sensors.4,6 For instance,
LCEs are being successfully integrated in robotics, micro-
electromechanical systems (MEMS),7 systems mimicking
muscle performances,8 and optical grating devices.9 Micro-
actuators have also specifically been developed via soft
lithography10 and microfluidic processes,11 while cantilevers
were obtained by inkjet printing.12
Received: June 4, 2014
Revised: July 15, 2014
Published: July 24, 2014
© 2014 American Chemical Society
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weakly cross-linked gel network by means of a prompt one-
dimensional stress. The latter induces a long-range orientation
of the mesogens and is kept until the completion of the second
reticulation step. Such stable, macroscopically oriented samples
are often referred to as “single-liquid crystal elastomers” or as
“monodomains”.15
propagation and reticulation to occur simultaneously. All these
new methods generally permit the preparation of macroscopi-
cally aligned thin elastomeric films as well as improve their
processing by novel techniques such as electrospinning,25
replica molding,27 and the enhancement of the anisotropic
factor via chemical approaches,31 swelling-induced electro-
optical effect,32 and piezoelectric and mechanical effects.24,31
Elastomers represent also an interesting class of stimuli-
responsive materials as they are intrinsically shape-memory
materials due to the hindered chain’ reptation33,34 mandated by
the reticulation. Essentially two types of shape memory effects
have been described in the literature for polymers. The one-way
type situation is commonly observed in amorphous polymers.5
Such a mode consists in the recovery of the initial “primary”
shape on heating, starting from an elongated “secondary” shape,
beforehand obtained at low-temperature upon a mechanical
stress. Lowering the temperature again, the primary shape is
conserved until the application of a subsequent mechanical
stress. This type is quite different from the second-way memory
shape effect, which consists in the continuous and reversible
shape-shift between two thermodynamically stable states. Both
types of shape memory effects occur in LCEs,5 but the second
mode is more specific for nematic LCEs,35 since very different
elongation amplitudes are generated by the large variations of
the order parameter with temperature (the cross-linking
hinders thermal motions of the mesogens but hardly the
thermal fluctuations of their main common orientation). This
variation is continuous and reversible36 between the most
elongated status close to the bottom limit of the N domain and
the maximum shrinking status, in the paranematic state, above
the smeared off transition (vide infra).5
Current efforts in LCEs focus essentially on (i) developing
performing synthetic methods (i.e., facile, reproducible,
adaptable), (ii) designing novel macromolecular network
topologies,3,4 and (iii) understanding and fostering predictions
of their self-organizing3,4 and thermoelastic behaviors to define
efficient construction strategies for the optimization of these
properties.16,17 Among the great diversity of all the possible
network structures realized so far,3,4 the so-called main-chain
LCEs (MC-LCEs)3,4 have recently generated a lot of interest.
Such a macromolecular network is achieved by the weak cross-
linking of linear polymeric chains, whose anisotropic pro-
mesogenic moieties are inserted in the polymer backbone.
Because of the intrinsic linear topology, these LCEs exhibit
extraordinary elastic properties, such as high conformational
anisotropy, high orientability, and large strain amplitudes,
which arise precisely from the direct coupling between the
long-range liquid-crystalline order (from the mesogens) and the
polymer chain network conformation.3,4 Main-chain systems
are therefore extremely appealing when strong deformations,
long cyclic actions, and high sensitivity to thermal stimuli are
required.
From a more fundamental point of view, the systematic
development of a wide range of MC-LCEs, integrating an
unrestricted variety of pro-mesogens, spacers, and cross-linkers,
is propitious to the better understanding of their intimate
structure−activity relationships as well as essential to the
investigation of theoretical models of their thermal-induced
elastic behaviors in the mesophases and in the isotropic
phase.16,18 In this respect, comblike LCEs (with mesogenic side
groups pending from the linear chains) have been widely
studied due to their relatively straightforward and accessible
syntheses.3,4 In contrast, such investigations are however only
emerging for MC-LCEs,4 consequently to numerous synthetic
difficulties encountered during their preparation (e.g., arduous
isolation of precursory linear polymers, low coupling, and cross-
linking efficiencies as reflected by high soluble contents, low
control of the network density, and low synthetic reproduci-
bility) and inadequate physicochemical properties (low
solubility, high transition temperatures, low thermal stability,
high propensity to form stiff solids, and rigid polydomain
materials), prejudicial to conduct reproducible measure-
ments.19−24 Aiming at improving the quality and accessibility
of MC-LCEs, several elegant synthetic approaches are currently
being explored, which include the photo-cross-linking of main-
chain prepolymers (laterally functionalized by photo-cross-
linkable units25 or with the photo-cross-linkable units inserted
within the polymer backbone26), the photoinduced addition of
thiols on olefins (aka thiol−ene photopolymerization) of
monomers containing olefinic and mercapto end-groups,27 or
the thermally activated cross-linking of end-unsaturated
polymers.28 Also, and of specific interest in this study, is the
“one-pot” approach,29 solely based on the chemistry of linear
siloxane-containing LC polymers.30 This highly versatile
synthetic methodology (vide supra) consists of the hydro-
silylation of divinylic liquid-crystalline monomers with linear
and cyclic siloxane comonomers, as chain-extender and
reticulating group, respectively, allowing therefore both chain
RESULTS AND DISCUSSION
■
Design. In this competitive and stimulating context, we
therefore endeavored the synthesis and characterization of two
series of weakly cross-linked mesomorphous MC-co-elastomers
showing a N phase by employing a combination of various
chemical approaches, including (i) the “one-pot” hydro-
silylation of divinylic liquid-crystalline monomers with linear
and cyclic siloxane comonomers, as chain-extender and
reticulating group, respectively;29 (ii) the introduction of
precise binary mixtures of mesogenic monomers, [N1x
+
N21−x], N1 showing individually a nematic phase (N) and N2 a
smectic C phase (SmC) over different temperature ranges, in
order to reduce mesophase symmetries and to lower transition
temperatures; (iii) the insertion of the siloxane fragments of
different length between mesogens, TMDS or HTMS, to yield
soft materials, with low transition temperatures; (iv) the
synthesis of mesogenic monomers bearing end-vinyl−ether
functional groups to enhance reactivity and thus to reduce
soluble contents;37,38 and finally (v) the use of a 5-fold cross-
linker, HD5, to ensure the formation of the network.31 Such a
design was elaborated in order to accurately investigate the
relationships between the mesophase structures and the shape
shifting and memory-shape behaviors of the elastomers.
The versatility of this methodology has been largely
demonstrated since it was first reported, and a large number
of siloxane-based MC-LCEs have been prepared accordingly,
merely requiring the preparation of customized divinylic-
terminated mesogens.29,31 This straightforward approach in-
deed presents several advantages: (i) it permits the readily
achievement of thermally stable MC-LCEs with low phase-
transition temperatures due to the regular alternation of
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a
Scheme 1. Synthetic Routes for the Monomers N1 and N2 (Intermediate Yields)
a
The detailed synthetic procedures are given in the Supporting Information: (i) DMAP, EDCI·HCl, CH2Cl2, 0 °C, 5−12 h; (ii) Pd/C (10%), H2,
THF, rt, 5−6 h; (iii) DIAD, PPh3, THF, 0 °C for 1 h and rt for 8 h.
siloxane and mesogenic units within the polymeric backbone,
(ii) offers the access to a wide range of macromolecular
structures consequently to the incorporation of the individual
constituents with unrestricted molecular shapes, and (iii) allows
straight and subtle modifications of the MC-LCEs structures by
easy permutations of the various components and modifications
of their respective proportions. Such flexibility in the choice of
the individual components is very unique and should in
principle permit complete systematic structure−activity inves-
tigations to tune the desired ultimate physical properties. A
recent improvement of the synthesis has just consisted in
considering LC monomers bearing vinyl−ether functional end-
groups, which considerably enhances the efficiency of the
hydrosylilation addition, as evidenced by lower soluble contents
determined in the corresponding MC-LCEs.37,38
The self-organizing and thermal behaviors of the two sets of
weakly cross-linked LCEs and of their related linear polymers
have been comprehensively examined by the combination of
POM (polarized-light optical microscopy), DSC (differential
scanning calorimetry), and SAXS (small-angle X-ray scattering)
techniques. Preliminary results of the thermoelastic behavior of
four representative LCEs have also been briefly described and
their elastic performances tentatively rationalized with the
mesophases’ structures.
presence of a lateral methyl group in one case (Scheme 1). The
synthetic route used for their preparation was performed in
three successive high-yield steps (Scheme 1, detailed
procedures in Supporting Information), starting with the
esterification of methylhydroquinone and 4-benzyloxyphenol
with 4-(benzyloxy)benzoic acid (Scheme 1, 1 and 3,
respectively), followed by the removal of the protecting benzyl
pending groups via catalytic hydrogenation (Scheme 1, 2 and
4) and by the subsequent binding of 4-vinyloxybutan-1-ol
fragments to the rigid cores 2 and 4 by the standard Mitsunobu
etherification. The overall yields after purification were 50 and
65% for N1 and N2, respectively. Three binary mixtures [N1x +
N21−x] were prepared with molar ratios x:(1 − x) = 0.25:0.75,
0.50:0.50, and 0.75:0.25 by dissolving the appropriate quantities
of each monomer in dry dichloromethane, followed by slow
solvent evaporation and vacuum drying (Supporting Informa-
tion).
As clearly identified by POM (Figure S4, Supporting
Information), N1 presents a N phase over a broad temperature
range (Cr 82 N 168 I),37 while the shorter N2 presents a SmC
phase over a small temperature interval (Cr 72 SmC 87 I).38
The long mesogen N1 logically leads to the broader liquid-
crystalline range (N), whereas the smectic phase vanishes in the
presence of the methyl group while perturbing the lateral
packing in layers.39 These two compounds are not miscible in
the crystalline state and likely form eutectic systems, but after
several heating−cooling cycles, the various mixtures behave
Monomers and Mixtures. In this study, two commonly
used mesogenic monomers were considered, N137 and N2,38
that principally differentiate by their core lengths and by the
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homogeneously as indicated by reproducible DSC traces
(Figure S8, Supporting Information). The smectic phase
already vanishes with the lowest N1 added concentration,
and all the mixtures show a broad N range, induced at
intermediate compositions by the melting temperature
depression and the linear variation of the clearing temperatures
(Figure 1 and Table S10, Supporting Information). Among the
polymeric chains. Moreover, to compare polymers and
elastomers, the polyaddition reactions were performed in the
same highly concentrated conditions as those used for the
synthesis of the elastomers described in the next section. They
were obtained by platinum-catalyzed hydrosilylation in toluene
between equivalent proportions of siloxane comonomers
(TMDS or HMTS) as chain spacers and the appropriate
mixtures of mesogens [N1x + N21−x] in different proportions
(Scheme 2, thereafter the polymers will be abbreviated as [N1x
+ N21−x+TMDS]p and [N1x + N21−x+HMTS]p, respectively;
Tables S1 and S2, Supporting Information).
All polymers and copolymers were purified via coprecipita-
tion (THF/MeOH) and the molecular weights analyzed by gel
permeation chromatography (GPC, Tables S3 and S4,
Supporting Information). Average degrees of polymerization
ranging from 40 to 245 units were obtained, with acceptable to
high polydispersity indices (PDI), comprised between 2.0 and
7.0 (PDI = Mw/Mn), most of them in the expected range for
such polyaddition reactions;30 the best results were found for
1
the TMDS systems. H NMR helped to confirm the retention
of the initial ratios of monomers after polymerization, by
comparing the total Si−CH2 integration with that of the methyl
signals from the N1 moiety (Table S5 and Figure S1,
Supporting Information).40
The structural features, common to all synthesized polymers,
are both (i) the presence of fairly long decoupling spacers
between consecutive mesogens conferring an alternated
polymeric structure30 and (ii) the existence of constitutional
regio-inhomogeneity caused by the random orientations and
distributions of the two mesogenic groups along the polymer
backbone, beneficial for the decrease of the lateral interactions
and thus to foster low transition temperatures (both monomers
are noncentrosymmetrical due to the presence of the lateral
methyl group in the central ring of N1 and to the polar ester
bond between the two aromatic rings in N2). As inferred by
POM (Figures S5 and S6, Supporting Information), DSC
(Figure S9, Figure S10, and Table S11, Supporting
Information), and SAXS (Figure S13, Supporting Information),
all the linear polymers exhibit thermotropic liquid-crystalline
behavior from room temperature up to ca. 100 °C. The overall
phase behavior of both series of polymers can be summarized as
displayed in the diagram below (Figure 2).
Figure 1. Comparative phase diagram of the monomers N1 and N2
and their three binary mixtures [N1x + N21−x] (x = 0.25, 0.50, and
0.75) as a function temperature. N, nematic phase; SmC, smectic C
phase; Cr, crystalline phases (Table S10, Supporting Information).
binary mixtures, the largest N temperature range is observed for
the equimolar mixture [N10.50 + N20.50] also having the lowest
melting transition temperature.
Main-Chain (Co)polymers. Two series of homogeneous
main-chain liquid-crystalline copolymers, alternating siloxane
fragments of different length, i.e. 1,1,3,3-tetramethyldisiloxane
(TMDS) or 1,1,3,3,5,5-hexamethyltrisiloxane (HMTS), with
mesogens N1 and N2, were then prepared (Scheme 2). In this
study, the polymers need to be synthesized with high
polymerization degrees (p), as they are the reference materials
of elastomers to be potentially compatible with actuating
applications, thus based on the interconnection of long
All the N1-containing polymers of the TMDS series exhibit
two mesophases, assigned to as SmC and N phases; the N2-
homopolymer, previously reported elsewhere, is deprived of the
N phase but exhibits instead an additional unidentified phase
below the SmC phase.38 Apart from the transition peaks, the
DSC traces of the TMDS copolymers (Figure S9, Supporting
Scheme 2. Sketched Illustration of the Synthesis of the Two
Series of Linear Polymers with Different Mesogenic
Monomer Ratios and Oligosiloxane Spacers ([N1x + N21−x
+
a
TMDS]p and [N1x + N21−x + HMTS]p)
Information) also contain a specific heat jump around −20
5
°C, attributed to the glass transition (Tg). While Tg is rather
insensitive to the binary composition, the clearing temperature
steadily increases with the N1 content, at a comparable rate as
in the monomeric mixtures, but with a lowering of about 60 °C.
The main difference with respect to the monomers and
mixtures consists in the induction of a SmC phase throughout
the series, below a transition temperature, also quasi-invariant
with composition. This phase is then retained down to room
temperature, replacing the crystalline range of the monomers.
In contrast, the N phase has totally vanished in the polymers
of the HMTS series, for which a direct transition from the SmC
phase to the isotropic state is observed (Figure S10, Supporting
Information). Clearing temperatures lie between the SmC-N
and the N−I phases’ transitions of the TMDS polymers,
a
Reaction conditions: toluene, PtCl2COD, 60 °C, 24 h (see details in
Supporting Information).
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Table 1. Structural Data of the Disiloxane (SmC and N) and
of the Trisiloxane (SmC) Polymers and Copolymers, [N1x +
a
N21−x + TMDS]p and [N1x + N21−x + HMTS]p
polymers
mesophase
d
hsi
hmes + hch
[N10.25 + N20.75
TMDS]p
+
+
+
SmC
N
27.6
6.5
4.5
ca. 30
ca. 5.0−5.5
ca. 5.0−5.5
[N10.50 + N20.50
TMDS]p
SmC
N
27.6
ca. 30
6.6
ca. 5.0−5.5
4.5
ca. 5.0−5.5
[N10.75 + N20.25
TMDS]p
SmC
N
28.5
ca. 30
6.6
ca. 5.0−5.5
4.7
ca. 5.0−5.5
[N2 + HMTS]p
SmC
SmC
28.3
28.6
6.5
6.5
4.6
4.6
[N10.25 + N20.75
HMTS]p
+
+
+
[N10.50 + N20.50
HMTS]p
SmC
SmC
SmC
28.5
29.0
29.7
6.8
6.6
6.6
4.7
4.7
4.8
Figure 2. Comparative phase diagram of both series of copolymers,
[N1x + N21−x + HMTS]p and [N1x + N21−x + TMDS]p, as a function
of the N1 content (mol %) and temperature. N, nematic phase; SmC,
smectic C phase; g, glassy state; X, unidentified phase. TMDS
homopolymers (x = 0, 1) were not made in this work, but transition
temperatures were taken from literature for comparison (refs 37 and
38).
[N10.75 + N20.25
HMTS]p
[N1 + HMTS]p
a
Abbreviations: SmC, smectic C mesophase; N, nematic phase.
Structural data measured at 25 °C for SmC phase and at 50, 64, and 68
°C in the N phase (with increasing N1 content): d, hsi, hmes, hch: layer
spacing and scattering maxima for lateral distances between siloxane
fragments, mesogenic monomers, and aliphatic chains, respectively (all
distances are given in Å).
therefore increasing with the N1 content (Figure 2). Glass
transitions are observed for two copolymers only, namely
[N10.50 + N20.50 + HMTS]p and [N10.75 + N20.25 + HMTS]p
(Tg = −20 and −13 °C, respectively); the other terms of the
series exhibit an additional phase, slightly below room
temperature, which could not be further investigated (thereafter
referred to as X).
SAXS patterns confirm in all cases the SmC structure of the
mesophase: they contain diffuse halos in the wide-angle region,
occurring from liquidlike lateral distances within sublayers, and
a sharp reflection in the small-angle region with a periodicity
compatible with tilted layers (Table 1 and Figure S13,
Supporting Information).
of the nanosegregation of the siloxane moieties.30,41 The
intercalation of the siloxane sublayers logically guarantees the
formation of overall long-range correlated layers but also
reduces interactions between successive mesogen-containing
strata and thus leads to the early vanishing of the long-range
correlated orientational order above the smectic range. The
impact on the mesomorphism is straightforward and less
important for the shorter spacer (TMDS series) and still
competes with the deleterious effect of the lateral methyl group
borne by the N1 mesogen (vide infra in TMDS elastomers),
resulting in their shrunken N ranges. In the HMTS series, the
mesomorphism is clearly dominated by the smectic phase and
the nanosegregation of the siloxane spacers.
Main-Chain (Co)-Elastomers. The synthesis of the main-
chain co-elastomers (Scheme 3) was performed according to
the previously described one-pot synthesis29,31 in a custom-
build centrifuge (Figures S17 and S19, Supporting Informa-
tion). The following protocol was kept identical for all systems:
the binary mixture of mesogenic monomers (N1x + N21−x), the
spacer (HMTS or TMDS), and the 5-fold cross-linker
(2,4,6,8,10-pentamethyl-1,3,5,7,9,2,4,6,8,10-pentaoxapentasile-
cane, HD5) were combined in a small 5 mL vial in adequate
proportions (Tables S6 and S7, Supporting Information),
dissolved in ca. 0.9−1 mL of freshly distilled toluene, and the
mixture heated gently to 60 °C until complete dissolution; in all
cases, the proportion of cross-linker HD5 was kept constant
throughout the entire series of compounds, i.e., 4−5 mol %
(e.g., 0.1 equiv/mmol). The catalyst was then added, and the
reaction mixture introduced in the customized cell, which was
then hermetically sealed. The mixture was left to react for 10
min, before centrifugation was started (60 °C, 1−3 h, 3000
rpm).
The halo at around 6.3−6.6 Å was readily assigned to the
siloxane sublayers (hsi), while the maximum at 4.6−4.8 Å
overlaps the scattering signals from distances between molten
aliphatic chains (hch) and mesogens (hmes), respectively. The
SmC periodicities lie in the same range and only vary little from
one series to another, e.g., 28 0.5 Å within the TMDS series
and 29 0.7 Å within the HMTS series. The patterns recorded
in the N phase of the N1-containing TMDS polymers are
similar to the ones in the smectic phase, apart from the trivial
broadening of the small-angle reflection. The correlation length,
associated with this broad feature, lies in the range of about 3
nm (Scherrer formula); that is, it extends to first neighbor’s
interactions. In the wide-angle region, the patterns contain an
unresolved broad scattering with a maximum around 5.0−5.5 Å,
overlapping the signals of hsi, hch, and hmes
.
Thus, as anticipated, all the alternated polymers resulting
from the combination of siloxane and mesogenic moieties are
mesomorphic, exhibiting SmC or SmC and N phases.30
Although their clearing temperatures are considerably below
the corresponding monomeric mixtures, the introduction of the
flexible siloxane segments widens the mesomorphic ranges to
room temperature. The siloxane spacers moreover induce the
formation of the smectic phase, systematically present, and
extending to higher temperatures with the longer spacer, while
the N range shrinks in TMDS polymers or vanishes totally in
HTMS polymers. This trend is just the inevitable consequence
Well-aligned elastomeric thin films were prepared via the
standard two-step method due to the flexibility introduced by
the siloxane segments (Figure S2, Supporting Information).29,31
The resulting fragile strips of gels were carefully removed from
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toluene infiltration). Surprisingly, the homoelastomers of the
HMTS series do not follow this expected behavior and swell in
an almost “isotropic” way.
Scheme 3. Sketched Illustration of the Synthesis of the Two
Series of Main-Chain Co-Elastomers ([N1x + N21−x + TMDS
+HD5] and [N1x + N21−x + HMTS + HD5]) with Different
a
Mesogenic Monomer Ratios and Oligosiloxane Spacers
The thermal behavior of both series of elastomers was mainly
characterized by DSC and SAXS; POM studies revealed
changes in the luminosity of the samples, but without
characteristic textures, providing information about the nature
of the mesophases formed, but not their complete elucidation.
Nevertheless, the high initial birefringence suddenly vanishes at
the temperature of the highest DSC peak, when the elastomers
transform into the isotropic-like state (Figure S7, Supporting
Information). Overall, the trend observed for the polymers
(Figure 2) can be translated to the corresponding elastomers
(Figure 3), evidencing that a low cross-linking density does not
affect greatly the thermal behavior.
a
Reaction conditions: toluene, PtCl2COD, 60 °C, 1−3 h, 3000 rpm.
the centrifuge cell, previously cooled down to room temper-
ature to lower reaction rate, and various rectangular lumps of
the elastomer were hung at one extremity. One-dimensional
alignment, e.g. parallel to the long strip axis, was achieved by
loading the other end with weights (a few milligrams) until the
liquid crystalline gel changed from turbid to transparent. The
films were then dried at 60 °C for 3 days and cooled down to
room temperature while still strained with the load. The
removal of the load lefts the elongated shape of the material
unchanged, likely because of the presence of the room
temperature LC phase (SmC). As evidenced previously,31 the
deformation from the initial shape preceding the orientation
process can be frozen into a “secondary” shape by crossing the
transition toward the smectic phase. Remarkably, this process is
purely due to mesomorphism and occurs for elastomers with
subambient glass transitions and therefore soft textures (like the
compounds studied here), in contrast to most of the shape-
memory polymer networks of the other types.5a The secondary
shape of smectic elastomers may remain unchanged after
several years aging.37
The completion of the alignment procedure consists in
washing the elastomers in a suitable solvent and extracting the
soluble content (sol %, corresponding to soluble part
embedded in the network) by means of successive Soxhlet
extractions in cyclohexane. A high soluble content yields poor
cross-linked, mechanically fragile networks. Hence, the lower
the soluble content, the better the quality of the material. Here,
the soluble contents determined were rather low, varying from
25 to 3% (Figure S3, Supporting Information), and are
comparable to the previously best reported siloxane-containing
elastomers.31,32 Swelling experiments were carried out in
toluene at 25 °C (Figure S3, Table S8, and Table S9,
Supporting Information). The co-elastomer samples presented
different dimensional changes in parallel (α∥) or perpendicular
(α⊥) with the direction of the applied stress during the
orientation process. The alignment of the cross-linked polymer
chains should reduce the swelling along the stress direction, in
accordance with high swelling anisotropies found for the
elastomers of both series (qz = α⊥/α∥, the swelling anisotropy,
is the ratio between the film dimensions before and after
Figure 3. Comparative phase diagram behavior of both series of co-
elastomers ([N1x + N21−x + HMTS + HD5] and [N1x + N21−x
+
TMDS + HD5]) as a function of the N1 content (mol %) and
temperature (°C). N, nematic phase; SmC, smectic C phase; X,
unidentified phase; g, glassy state. TMDS homoelastomers (x = 0, 1)
were not made in this work, but transition temperatures were taken
from literature for comparison (refs 37 and 38).
The elastomers of the TMDS series show the same sequence
of states as the corresponding polymers, except that the
transition temperatures are substantially increased: Tg is
increased by roughly 15 °C, the SmC to N transition by
about 20 °C, and the N to isotropic state (actually paranematic-
like state, IN, vide supra) by on average 25 °C (Figure S11 and
Table S12, Supporting Information). On the contrary, the
thermal behavior of the HMTS elastomers is almost the carbon
copy to that of the corresponding polymers: they show at
almost the same temperatures, the SmC phase, the glassy state,
or a subambient unidentified phase “X” (Figure S12 and Table
S12, Supporting Information). The only significant difference
consists in the appearance of an additional mesophase above
the SmC phase, likely a N phase, for the materials with the
highest N1 contents (for x ≥ 0.5). Although the mesophase
temperature range is too narrow (a few degrees) for a definitive
confirmation by SAXS, the reappearance of the N phase could
be explained by the slightly thinner siloxane sublayers as
compared to that of the polymers and also possibly to the
partial disordering of the smectic layers upon the introduction
of the disturbing cross-linker (vide infra).
SAXS patterns of all co-elastomers are consistent with
aligned SmC materials, as they contain diffuse bows centered
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on the equatorial plane for the scattering maxima hSi, hch, and
h
mes and a set of four small-angle Bragg spots revealing the cone
distribution of the lamellar first-order reflection around the
elongation direction, parallel to the meridian (Figures 4 and 5,
Figure 5. X-ray data of co-elastomer [N10.50 + N20.50 + HMTS +
HD5], chosen as a representative example. Top: aligned SAXS pattern
in the SmC phase at 25 °C (the direction of elongation is parallel to
the meridian); middle: radial profiles over 360°; bottom: azimuthal
profiles in the wide- and small-angle regions (the direction of
elongation corresponds to ψ = 0, 180°).
When the isotropic state (or paranematic-like, vide supra) is
reached, the small- and wide-angle scatterings blur out the
broad halos. Residual orientations can however be detected for
the TMDS elastomers only, but this peculiar behavior only
comes in the forefront on cooling in the mesophase: the
elastomers of this series reproduce similar oriented patterns
than on heating, while the elastomers of the HMTS series give
rise to patterns with continuous rings, probing the isotropically
oriented smectic domains. A deeper insight in the recovery
process of the TMDS elastomers was deduced from SAXS
patterns of [N1x + N21−x + TMDS + HD5] by scrutinizing the
variations of ξ and S (Figure 6 and Table 2). In the SmC phase,
the evolution of ξ trivially agrees with long-range correlations
(25−30 nm), but the transition to the N phase only causes a
drop to about half its length, preceding a gradual decrease in
the N domain. This differs from the classical behavior of
calamitic mesogens,42 with the immediate collapse of ξ to short
range at the transition to the N phase. The variation of S is even
more unconventional: instead, as anticipated, of a continuous
decrease in the N phase, and the final vanishing at the transition
to the isotropic state, S remains constantly large in the N
domain and progressively decreases close to and above the
transition to the isotropic state. Measurable orientations are still
observed at higher temperatures, far beyond the transition to
the isotropic state, an indication that the material is maintained
in some sort of intermediate transitory state, clearly not a
kinetic but a real thermodynamic state, between the nematic
Figure 4. X-ray data of co-elastomer [N10.50 + N20.50 + TMDS +
HD5], chosen as a representative example. Top: aligned SAXS
patterns in the SmC phase at 25 °C, left, and in the N phase at 68 °C,
right (the direction of elongation is parallel to the meridian); middle:
radial profiles over 360° of the corresponding mesophases (d001, hsi,
hmes + hch: first-order lamellar periodicity, and scattering maxima for
lateral distances between siloxane units, monomers, and chains;
bottom: azimuthal profiles of both mesophase in the wide- and small-
angle regions (the direction of elongation corresponds to ψ = 0,
180°).
Table 2, Figures S14 and S15, Supporting Information). The
layer spacings of the TMDS (25.5 1.0 Å) and HMTS (28
1.0 Å) elastomers SmC phases are slightly reduced with respect
to the periodicities of the SmC phases of the corresponding
polymers (by about 2 and 1 Å, respectively); therefore, the
mesogens are slightly more tilted, mainly due to the expansion
of the cross-linkers within the siloxane sublayers. When TMDS
elastomers are heated in the N phase, the only important
change in the SAXS patterns consists in the widening of the
small-angle spots, the alignment remaining almost the same.
This widening is however limited, and corresponds to
correlation lengths of about 4 times those found for the
TMDS polymers, amounting to clusters of 4−5 neighboring
layers. Comparable large positional correlations have been
previously found in the N phase of other related main-chain
elastomers.38
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Table 2. Structural Data of the Disiloxane (SmC and N) and Trisiloxane (SmC) Co-Elastomers, [N1x + N21−x + TMDS + HD5]
a
and [N1x + N21−x + HMTS + HD5] (4 mol % HD5)
elastomers
mesophase
d
ψchev
ξ [n]
S
hsi
hmes + hch
[N10.25 + N20.75 + TMDS + HD5]
SmC
N
25.5
27
41.5
49
30 [12]
11 [4]
0.90
0.89
0.91
0.80
0.91
0.90
0.83
0.89
0.75
0.80
n.d.
6.5
n.d
6.6
n.d.
6.8
n.d.
6.6
6.8
6.8
6.7
6.7
4.5
4.5
4.4
4.6
4.5
4.5
4.6
4.6
4.6
4.6
4.6
[N10.50 + N20.50 + TMDS + HD5]
[N10.75 + N20.25 + TMDS + HD5]
SmC
N
24.6
28
40
>60 [>25]
12 [4]
44
SmC
N
26.5
27
43
26 [10]
12 [4]
45
[N2 + HMTS + HD5]
SmC
SmC
SmC
SmC
SmC
26.9
28.1
28.5
28.0
28.6
30
30 [11]
24 [9]
[N10.25 + N20.75 + HMTS + HD5]
[N10.50 + N20.50 + HMTS + HD5]
[N10.75 + N20.25 + HMTS + HD5]
[N1 + HMTS + HD5]
32
26
34 [12]
29 [10]
29 [10]
37.5
39
a
Structural data measured at 25 °C, in the SmC phase and at T = 65, 68, and 85 °C, respectively, in the N phase: d, layer spacing (Å); ψchev, chevron
tilt angle (deg); ξ: correlation length (nm); n: number of correlated layers; S: nematic order parameter; hsi, hmes + hch: scattering maxima for lateral
distances between siloxane units, monomers and chains, respectively (Å). n.d.: not determined due to weak signal.
To understand further the relationship with the siloxane
spacer, the nanosegregation process needs to be more deeply
investigated by accessing to the bulkiness of the different
species, especially their partial volumes and cross sections,
obtained from reference dilatometric measurements. The
individual cross section of mesogens, σmes, aliphatic chains,
σch, and siloxane units, σsi, can then be compared to the
molecular area Amol, assimilated to the statistical layer portion
covered by a single repeat unit, defined as ratio between the
repeat unit volume, Vmol, and the measured layer spacing, d.44
The main discrepancy in these systems lies between the cross
sections of siloxane (σsi ≈ 45 Å2) and of the other constitutive
moieties (σch ≈ 21.3 Å2, σmesN2 ≈ 23.6 Å2, and σmesN1 ≈ 26.5
Å2). The effective molecular areas of the polymers and
elastomers take intermediate values, ranging from 32 to 40
Å2, and involve several compensation mechanisms. The
Figure 6. Temperature variations of the order parameter, S, and of the
accommodation of the bulky siloxane segments is possible
through the ruffling of their sublayers,45 which occurs more
easily with thinner sublayers and hampered by the network, in
consistency with the experimental variation of Amol between
both series (Table 3). Molecular areas remain however close to
siloxane, as the other sublayers more easily expand, through the
folding of alkyl chains and the tilting of mesogens.
correlation length of smectic layers, ξ, for the co-elastomer [N10.75
N20.25 + TMDS + HD5], chosen as a representative case. IN:
paranematic-like state state; SmC + N, N + IN: coexistence regions.
+
and the isotropic liquid, referred to as paranematic-like state.
Clarifying the real nature of such a state would require
extensive phase transition studies far beyond the scope of this
investigation.43 Thereafter, the term “paranematic-like” state
will be used and labeled as IN.
The residual orientational order is obviously the main reason
for the recovery of the oriented SAXS patterns on cooling and
for the delay of the transitions to higher temperatures as in the
polymers. The trend observed with the HMTS elastomers
contrasts sharply with this thermal behavior, as their transitions
to the isotropic state are not delayed and erase definitively the
orientation of the smectic domains. This behavior would a
priori be consistent with classical transitions to isotropic states,
thereafter labeled “I”.
Therefore, the reticulation of the polymeric chains only
modifies slightly the polymorphism: the HMTS elastomers
mainly display a SmC phase, over almost identical temperature
ranges as the nonreticulated systems, and both N and SmC
phases are preserved in the TMDS systems. However, the
network drastically affects the phase transition process, with
very different behaviors within both series: reversible changes
preserving the memory of the secondary orientation for the
TMDS systems and irreversible tendencies to return to the
initial unoriented state for the HMTS systems.
The ratio σmes/Amol gives access to the tilt angle, ψmes
,
representing the average polar tilt angle of the mesogen axis
from the layer normal. This information is completed by the
chevron angle, ψchev, which is directly measured on the oriented
SAXS patterns and corresponds to the cone distribution of the
smectic layers normal or to the piling direction in the smectic
correlation volumes of the nematic phase (Figure 7). The cone
axis coincides with the direction of elongation and therefore
with the preferential direction of the axes of the polymer chains
and of the mesogens. Both tilt angles may be different since
they do not relate at the same scale: ψchev is the macroscopically
emerging tilt and possibly integrates an azimuthal distribution
of mesogens’ axes, while ψmes refers to the local packing within
layers and thus sets the upper limit of ψchev. The sharpness of
the small-angle spots and the closeness of both tilt angles in the
TMDS elastomers (Figure 4 and Table 3) confirm the uniform
mesogen alignment. The tilting is moreover preserved in the N
phase within smectic-like clusters (Table 2), the slightly higher
angles having different interpretations (thermal expansion,
small azimuthal component, etc.). The HMTS elastomers show
on the contrary a clear gap from ψchev to ψmes, without the
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Table 3. Structural Data of Polymers and Elastomers in the
a
SmC Phase at 25 °C
polymers/
elastomers
Vmol
ψchev
Amol
σmes
ψmes
(deg) (deg)
Δψ
(Å3)
d (Å) (deg) (Å2)
(Å2)
[N2 + TMDS]p
849
894
23.6
24.3
[N10.25 + N20.75
TMDS]p
+
+
+
27.6
27.6
28.5
32.4
34.0
34.5
41
[N10.50 + N20.50
TMDS]p
939
984
25.1
25.8
42.5
42
[N10.75 + N20.25
TMDS]p
[N1 + TMDS]p
1029
26.5
23.6
[N2 + TMDS +
HD5]
843*
[N10.25 + N20.75
+
889
25.5
24.6
26.5
42
40
43
34.9
38
24.3
25.1
25.8
26.5
46
49
46
4
9
3
TMDS + HD5]
[N10.50 + N20.50
TMDS + HD5]
[N10.75 + N20.25
+
934
+
979
36.9
TMDS + HD5]
[N1 + TMDS +
HD5]
1023*
[N2 + HMTS]p
974
28.3
28.6
34.4
35.6
23.6
24.3
47
47
[N10.25 + N20.75
HMTS]p
+
+
+
1019
[N10.50 + N20.50
HMTS]p
1064
1109
28.5
29.0
37.3
38.3
25.1
25.8
48
48
[N10.75 + N20.25
HMTS]p
[N1 + HMTS]p
1154
957
29.7
26.9
38.9
35.6
26.5
23.6
47
[N2 + HMTS +
HD5]
30
32
26
38
39
48.5
18.5
15
21
21
9
[N10.25 + N20.75
+
1002
1047
1092
1137
28.1
28.5
28.0
28.6
35.7
36.7
39.0
39.8
24.3
25.1
25.8
26.5
47
47
49
48
HMTS + HD5]
[N10.50 + N20.50
HMTS + HD5]
[N10.75 + N20.25
+
+
HMTS + HD5]
[N1 + HMTS +
HD5]
a
Partial molecular volumes at 25 °C: Vmol is the volume of the
molecular motif, Vmol = xVN1 + (1 − x)VN2 + VHMTS/TMDS for the
polymers, Vmol = xVN1 + (1 − x)VN2 + 0.9VHMTS/TMDS + 0.04VHD5
(*Vmol = xVN1 + (1 − x)VN2 + 0.88VHMTS/TMDS + 0.05VHD5) for the
elastomers; VN1 = 813 Å3, VN2 = 633 Å3, VTMDS = 216 Å3 (polymers),
VTMDS = 195 Å3 (elastomers), VHMTS = 342 Å3 (polymers), VHMTS
=
307 Å3 (elastomers), VHD5 = 421 Å3. Other structural data: d, layer
spacing; Amol = Vmol/d, molecular area; σN1, σN2, cross sections of
mesogens N1 and N2, respectively; σmes = xσN1 + (1 − x)σN2, average
cross-section; ψmes = arcos(σmes/Amol), ψchev, Δψ = ψmes − ψchev
mesogen tilt angle, chevron tilt angle and difference.
,
Figure 7. Schematic representations of the molecular packing and
corresponding X-ray patterns, in the SmC phase of the stretched
elastomers with different lengths of siloxane spacers. Top: TMDS
spacer. Middle: HMTS spacer. Bottom: predictive structure with
broadening of the small-angle spots (Figure 5 and Table 3).
This means that the orientations of the mesogens deviate from
one another, but that the preferential orientation of the smectic
domains is still the direction of elongation. In other words,
smaller average tilt angles ψchev emerge from the contributions
of zones with different azimuthal component (Figure 7). The
orientation of the mesogens is poorly correlated in the HMTS
elastomers, and the trend between both series proves to be the
expression on the SmC structure of the same feature driving the
evolution of the polymorphism: the nanosegregation of the
siloxane sublayers isolating successive mesogen sublayers from
each other (vide infra). With the further lengthening of the
siloxane spacers (increase of the volume fraction), the structure
should predictably continue evolving in the same way; the
longer spacer. f is siloxane volume fraction; n and n represent the
⃗
⃗
si
0
directions of the mesogens’ axes and of the normal to the layers; ψmes
and ψchev are the mesogen and the chevron tilt angles; Amol is the
molecular area; σsi is the cross section of the siloxane segments.
Corollary, the influence of the siloxane sublayer thickness
also provides a molecular interpretation for the observed phase
behavior change: the thicker HMTS sublayers reduce the
correlations between the orientations of the mesogens from
successive layers and leave a classical isotropic state when the
smectic layers vanish. Conversely, above the broad N ranges of
the TMDS elastomers, the orientational long-range order is
macroscopic tilt might then totally vanish (ψchev = 0 with ψmes
>
45°) and yield De Vries-like smectic A phases (Figure 7).46
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replaced by residual orientational correlations at a level set by
the strength of the cross-linking.
The elastomeric network associates a given structure of
cross-linked polymer chains to a shape. The analyses of the
phase behavior and of the molecular packing were therefore
partially completed with the measurement of the shape
variation as a function of temperature for the four following
elastomers: [N1x + N21−x + TMDS + HD5] with x = 0.25 and
0.75 and [N1x + N21−x + HMTS + HD5] with x = 0 and 0.5.
These preliminary thermoelastic experiments were realized with
an elongation setup by starting from the “secondary” shape,
corresponding to the oriented state fixed during the second
preparation step. Measurements were conducted on both
heating and cooling cycles by changing the temperature
stepwise with 50 min stabilization time at each step.
Elongations were calculated in reference to the sample size at
the highest measured temperature beyond the mesomorphic
range. Data obtained on successive heating and cooling runs for
both TMDS elastomers superimpose fairly well in the SmC, N,
and IN ranges, except a small hysteresis effect around the N−
SmC transition due its first-order character (Figure S16,
Supporting Information). The superimposition confirms that
the recovery of the secondary orientation in SAXS patterns
(vide infra) combines with the recovery of the secondary shape.
Consequently, these elastomers exhibit definitively two-way
shape-memory effects, obviously in connection with the
residual orientational correlations that initiate the growth of
the low-temperature oriented states. As for the HMTS
elastomers, the thermoelastic experiments prove that the
definitive erasing of the secondary orientation in the isotropic
state underlies the shrinking to a permanent shape close to the
initial unoriented state. The lengthening of the siloxane spacer
therefore changes the shape-memory behavior for the one-way
type. Recall that the cross-linker content is the same in both
series of co-elastomers (ca. 1 unit for 25 divalent siloxane units,
Tables S6 and S7, Supporting Information), and therefore the
density of the cross-linking is unlikely to be at the origin of the
switching behavior between both series.
Figure 9. Thermoelastic first heating curves of HMTS elastomers,
[N1x + N21−x + HMTS + HD5] (x = 0.50, 1.00) as a function of the
temperature interval to the transition between nematic and isotropic
states (TI). The dotted line corresponds to the SmC−N transition of
[N1 + HMTS + HD5].
HMTS homoelastomer. These values are difficult to interpret
as they are influenced by various experimental parameters, such
as the exact sample dimensions modified by pinching zones, the
elastic properties after the first preparation step depending on
the degree of cross-linking, on the relative homogeneity of the
network and on the various chain conformations and topologies
(opened and closed loops, hairpins, knots, etc.), and on the
composition, the polymorphism, and the gap between
mesogens and chevron tilt angles, etc. Isolating each influence
would be an extensive experimental work far beyond the
present study. Whatever the initial elongation, the shape
remains constant in the SmC phase; the shrinking always starts
a few degrees below the transition to the next phase and ends a
few degrees above the transition to the isotropic state. Since the
homo-elastomer and the co-elastomer of the HMTS series
show respectively no and a tiny intermediate N phase, the
whole shrinkage takes place abruptly, within a narrow
temperature range. To the contrary, the shrinking of the
TMDS elastomers is gradual and spreads over the broad N
range. The rate of shrinking is nevertheless variable, from rather
moderate in the middle of the nematic domain, to steep around
and inside transitional zones, in connection with the variations
of S and ξ deduced from SAXS patterns (vide infra).
Apart from the reversibility of the shape change on (first)
heating, the elastomers differ from one another in their initial
elongations and in their rates of shrinking (Figures 8 and 9).
Initial elongations lie around 200% for both the TMDS co-
elastomers and the HMTS co-elastomer but reach 350% for the
To summarize, these prelimary thermoelastic experiments
confirm that the reversible X-ray pattern orientations of TMDS
elastomers effectively underlie the gradual and reversible shape
shift throughout the broad nematic range, whereas the once
oriented HMTS elastomers irreversibly shrink to the
unoriented shape as soon as the long-range correlated smectic
layers collapse. It should be emphasized that the shape recovery
of the TMDS elastomers was not affected by the time stayed in
the paranematic-like state (up to a day) and that no delay was
observed for the irreversible shrinking of the HTMS
elastomers. Both series of elastomers turned out to be classical
soft materials whose properties are determined by the
molecular organization.
CONCLUSION
■
A binary system, composed of mesogens with the classical
calamitic architecture (rodlike cores of different bulkiness and
terminal vinyl−alkyl tails), was incorporated in main-chain
polymers via direct hydrosylilation reaction with various linear
Figure 8. Thermoelastic heating curves of TMDS elastomers [N1x +
N21−x + TMDS + HD5], (x = 0.25, 0.75) as a function of the
temperature interval to the transition between nematic and para-
nematic-like states. SmC + N: coexistence temperature region.
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siloxane oligomers. Instead of the SmC phase shown by the
small monomer, the longer mesogen and all the homogeneous
mixtures give rise to broad nematic ranges, consequent to the
perturbations of the lateral packing (introduced by the methyl
side-group borne by the long mesogen). The nanosegregation
of the siloxane spacers confined in intermediate sublayers
however counteracts this effect in the resulting alternating
polymers, as proven by the emergence of the SmC phase, either
below a somewhat shrunken nematic range or with a direct
transition to the isotropic state. In addition, the siloxane spacers
impose tilt angles close to 45° due to the large difference
between the cross sections of both the siloxane fragments and
the mesogens.
The mesomorphic properties were marginally affected upon
the replacement of a small fraction of the siloxane spacers by
cross-linking segments to produce the elastomers. In particular,
the room-temperature SmC phase and the subambient glass
transition were preserved, which allowed the alignment realized
by mechanical elongation during a step of the synthesis and its
retention in the final soft material. This alignment provided a
deep insight into the intricate structure, especially through the
comparison of tilt angle of mesogens and chevron tilt angle of
the smectic planes, on the one hand, and through the access to
the orientational order parameter and to the smectic correlation
length, on the other hand. These detailed structural
investigations carried out by SAXS were further completed by
a short description of our preliminary thermoelastic experi-
ments, measuring the connected variations of shape.
of the dynamics of the entire network chains conformation on
the drastic shape shift.
Therefore, the results obtained in this study appear to be
relevant for the designs of shape-memory LC elastomers, with a
fairly good anticipation of their elastic and thermal properties.
It would also be of interest in future works to extend the variety
of spacers to be inserted in the main chain of the polymers to
other types of siloxane oligomers of various bulkiness and
intricate structures as well as to corresponding mixtures. The
use of controlled mixtures of mesogens also proved as an
efficient mean to considerably lower the transition temper-
atures with the partial retention of the mesomorphic behavior
of the initial components, and in some cases the survival of the
N mesophase. The drastic change between the elastic behaviors
(vanishing of the elongated shape memory and of the
paranematic-like state between both series) naturally puzzles
over the features in the switching zone at intermediate spacer
lengths and would merit specific investigations in future works.
ASSOCIATED CONTENT
* Supporting Information
■
S
Experimental part, tables of the thermal behavior of the
monomers, copolymers and co-elastomers, selected POM
images, DSC traces, SAXS patterns, thermoelastic curves, and
centrifuge pictures. This material is available free of charge via
the Internet at http://pubs.acs.org.
AUTHOR INFORMATION
Corresponding Author
*E-mail Bertrand.Donnio@ipcms.unistra.fr (B.D.).
■
The length of the siloxane spacer in particular turns out to be
an efficient and essential parameter to confer shape-memory
behavior of different types to the elastomers, i.e., of the one-way
type for the HMTS system, for which the elongated shape is
erased on heating and not recovered on cooling, and of the
two-way type for the TMDS system, for which the shape of the
samples reversibly changes between states, distinct from the
initial shape before the cross-linking step.
Notes
The authors declare no competing financial interest.
ACKNOWLEDGMENTS
■
A.G.M., B.H., N.B., D.G., and B.D. thank CONACyT
scholarship 185485 for financial funding during this research
work (A.G.M.), the European Union through the Research
Training Networks project FUNCTIONAL LIQUID-CRYS-
TALLINE ELASTOMERS (FULCE-HPRN-CT-2002-00169),
For the elastomers with the shorter siloxane spacer, relatively
broad nematic ranges were maintained, and strong correlations
between the orientations of the mesogens persisted above the
mesomorphic range, in consistency with the existence of the
paranematic-like state beyond the N phase and with the
vanishing gap between mesogens’ and chevron tilt angles.
These features translated into gradual and reversible variations
of elastomers shape with temperature, from constant
elongations of about 200% in the SmC phase to almost no
elongation in the paranematic-like state. This two-way shape-
memory behavior contrasted with the one-way type observed
for the other elastomer series with the thicker siloxane
sublayers, evidenced by the reduced chevron tilt angle and
predominant confinement in sublayers, thus interrupting more
efficiently the orientational correlations between mesogens of
successive layers. Consistently, the N phase disappeared or was
only preserved in a narrow range, while the initial elongations
and orientations vanished rapidly and irreversibly close to the
transition to the isotropic state. The modifications of the
mesogen’ orientation and of the sample shape during
preparation and thermal treatment necessarily also translate
into changes of the entire macromolecular network con-
formation. Fruitful extensions of this work could be scattering
experiments on elastomers with labeled mesogens, NMR
spectroscopy, and computational modeling taking into account
the mesogens segregation in order to better appreciate the role
́
CNRS, and the Universite de Strasbourg for financial support,
and Professor Dr. Dr. h.c. Heino Finkelmann for helping in the
thermoelastic apparatus and spin-casting centrifuge construc-
tions.
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