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Heat-Induced Actuator Fibers: Starch-Containing Biopolyamide
Composites for Functional Textiles
Hossein Baniasadi,
§
Zahra Madani,
§
Mithila Mohan, Maija Vaara, Sami Lipponen, Jaana Vapaavuori,*
and Jukka V. Seppälä*
Cite This: https://doi.org/10.1021/acsami.3c08774
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ABSTRACT: This study introduces the development of a thermally responsive
shape-morphing fabric using low-melting-point polyamide shape memory
actuators. To facilitate the blending of biomaterials, we report the synthesis
and characterization of a biopolyamide with a relatively low melting point.
Additionally, we present a straightforward and solvent-free method for the
compatibilization of starch particles with the synthesized biopolyamide, aiming to
enhance the sustainability of polyamide and customize the actuation temperature.
Subsequently, homogeneous dispersion of up to 70 wt % compatibilized starch
particles into the matrix is achieved. The resulting composites exhibit excellent
mechanical properties comparable to those reported for soft and tough materials,
making them well suited for textile integration. Furthermore, cyclic
thermomechanical tests were conducted to evaluate the shape memory and
shape recovery of both plain polyamide and composites. The results confirmed
their remarkable shape recovery properties. To demonstrate the potential application of biocomposites in textiles, a heat-responsive
fabric was created using thermoresponsive shape memory polymer actuators composed of a biocomposite containing 50 wt %
compatibilized starch. This fabric demonstrates the ability to repeatedly undergo significant heat-induced deformations by opening
and closing pores, thereby exposing hidden functionalities through heat stimulation. This innovative approach provides a convenient
pathway for designing heat-responsive textiles, adding value to state-of-the-art smart textiles.
KEYWORDS: compatibilization, copolyamide, starch, shape memory actuator, heat-responsive smart textile
■INTRODUCTION
Polyamides (PAs) have long been recognized as one of the
major engineering thermoplastics, finding applications in
various industries such as sports, furniture, textiles, automotive,
fishing, and medical devices. Their exceptional properties,
including excellent chemical resistance, electrical insulation,
robust mechanical and thermal characteristics, as well as ease
of processing, have contributed to their widespread use.
1−4
Recently, PA-based fishing lines have attracted attention for
their intriguing application as thermally responsive artificial
muscles capable of delivering significant strokes through shape
memory eects.
5,6
These materials oer direct compatibility
with textile manufacturing techniques, allowing seamless
integration into passive textiles and transforming them into
dynamically adaptive three-dimensional (3D) networks.
7
In
this study, we present the design of a predominantly biobased
low-melting-point polyamide, which opens up new possibilities
in lowering the actuation temperature of previously reported
PA-based artificial muscles. Simultaneously, it reduces the
reliance of smart textile materials on petroleum-based raw
materials. By exploring this approach, we aim to expand the
application areas of polyamides and promote sustainable
alternatives to the development of smart textiles.
Conventionally, PAs are synthesized from petroleum-based
monomers; nevertheless, accumulating more CO2into the
atmosphere and environmental concerns emerging from their
gradual degradation rates and harmful degradation products,
along with the continuous decline in oil reserves, have sparked
a growing interest in more environmentally friendly ways of
synthesizing PAs.
8,9
Accordingly, several partially and fully
biobased PAs like PA410, PA510, PA1010, and PA11 have
been developed recently from natural sources, e.g., castor oil,
endowing comparable mechanical strength to traditional
petroleum-based PAs like PA6 or PA66.
10,11
The other well-
established approach toward developing more sustainable
materials is the incorporation of conventional synthetic
polymers, e.g., PAs, with biomass. The production of these
biocomposites emits proportionally lower amounts of carbon
and greenhouse gases into the atmosphere than plain
Received: June 18, 2023
Accepted: September 21, 2023
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petroleum-based plastics; furthermore, replacing conventional
PAs with these biocomposites can reduce the usage of fossil
fuel reserves.
12−15
Nevertheless, compounding biomass, i.e.,
biofillers with PAs, is limited by the relatively low thermal
stability of biofillers. Moreover, due to the high polarity and
rich intermolecular and intramolecular hydrogen bonds in PAs,
high processing temperatures, causing decoloring and thermal
degradation of biofillers during blending, are needed.
16,17
Starch is one of the potential carbohydrate candidates for
blending with polymers thanks to its abundance, low cost,
inherent biodegradability, and renewability.
18,19
Nevertheless,
due to the abundant surface hydroxyl groups, starch suers
from low compatibility with most polymers, e.g., PAs.
20
Numerous approaches have been developed to enhance starch
compatibility with polymer matrices. The most well-known
and common method is employing polyethylene-grafted maleic
anhydride as a compatibilizer.
21,22
The grafting of oligomers/
polymers is another ecient strategy to enhance interfacial
adhesion between starch and polymer matrix.
23,24
Likewise,
disrupting hydrogen bonding between starch molecules via an
esterification reaction between hydroxyl groups of starch and
succinic anhydrides has also been of interest.
25
Accordingly,
alkenyl succinic anhydrides with dierent chain lengths, from
octenyl succinic anhydride to octadecenyl succinic anhydride,
could be used for starch treatment. The size of the alkenyl
group and the degree of substitution are important parameters
determining the level of starch hydrophobic. Octadecenyl
succinic anhydride (OSA) is the most popular alkenyl succinic
anhydride, which has been used for a long time for starch
hydrophobization. OSA-treated starch containing 3.0 wt %
Figure 1. Schematic representation of (a) copolymerization and (b) surface modification of starch with OSA molecules.
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OSA has even been approved by FDA in foods.
26−29
Besides, it
has been used as a sustainable template for biomedical
applications.
30
Although there are many reports on OSA-
treated starch as a stabilizer and emulsifier in many food
systems,
31,32
only a few reports are available on its melt
blending with polymer matrices.
26,33
In the current study, OSA was grafted on the surface of
starch particles (OSA-g-starch) through a simple and solvent-
free method to not only strengthen polymer/particle interfacial
adhesion but also improve starch dispersion in the matrix
significantly. Furthermore, to prevent any thermal degradation
of starch particles during compounding, for the first time, a
novel biobased polyamide with a remarkably low melting point
of 135 °C was synthesized through a copolymerization process.
Thanks to the employed compatibilization methods, various
concentrations of OSA-g-starch, including a very high
concentration of 70 wt %, were easily melt-blended with the
synthesized copolyamide at 160 °C. To the best of our
knowledge, there is no other comprehensive investigation of
biopolyamide/starch composites in the literature. Moreover, to
showcase the advantages of the prepared biocomposites in
shape memory textiles, coiled actuators were prepared and
integrated into a woven 3D textile structure. The demonstrator
fabric was able to switch between closed and open porous
structures as a function of temperature, thus opening the door
to the wide variety of multifunctional textiles, where an added
functionality could be exposed and hidden on command.
■EXPERIMENTAL SECTION
Materials. Starch from potato, 11-aminoundecanoic acid, sodium
hypophosphite monohydrate (>99%), trifluoroacetic anhydride
(reagentPlus, ≥99%), and chloroform-d(99.8 atom % D) were
obtained from Sigma-Aldrich. Octadecenyl succinic anhydride
(mixture of isomers) and 1,12-diaminododecane (≥98%) were
prepared from TCI, Japan. Chloroform (for analysis EMPARTA
ACS) was purchased from Merck. 1,18-Octadecanedioic acid was
bought from Cathay Biotech Company, China.
Copolymerization. The partially biobased polyamide in this
study was synthesized through a copolymerization reaction involving
two dierent polyamides, namely, PA11 and PA1218. The reaction
process is illustrated schematically in Figure 1a. The monomers,
including 11-aminoundecanoic acid, 1,12-diaminododecane, and
octadecanedioic acid, were introduced in equal molar quantities
into a stainless steel reactor equipped with a heating jacket and an
overhead mixer. Sodium hypophosphite monohydrate was added as a
catalyst, and the reactor was heated until the temperature reached 200
°C. The monomers were then melted for 1 h under a nitrogen
atmosphere at this temperature. Subsequently, the temperature was
increased to 240 °C, and the molten monomers/oligomers were
gently mixed for 4 h under a nitrogen stream to complete the
polycondensation reaction.
After the reaction, the reactor was cooled under a nitrogen flow,
and the resulting product, a copolymer of PA11 and PA1218 (termed
PA11coPA1218), was collected. The synthesized copolyamide was
Figure 2. (a) Schematic representation of the developed fabric. The top image shows the cross-sectional diagram, and the bottom image presents
the layer diagram constructed from the cross-sectional diagram to denote the placement and interlacement of each layer. In all layers, the warp is
100% cotton. In layer 1, the weft is 100% cotton, and the developed actuators, while in the rest of the layers, the weft is Linen-Tencell. (b) Digital
image from the top view of the fabric woven with PSMS50 actuators and cotton in the first layer. (c) Digital image of the fabricated smart textile.
(d) Image of fabric photographed in a dark chamber showing the eects of glow-in-the-dark yarn stripes interlacing.
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ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
C
milled using a Retsch SM 300 Cutting Mill with a 6 mm sieve size
(square holes). For comparison, the parent homopolymers, PA11 and
PA1218, were also synthesized by using the same polymerization
conditions as reference materials.
Compatibilization. The starch was surface-treated to be
compatible with the biopolyamide matrix. The treatment was carried
out by grafting octadecenyl succinic anhydride (OSA) via a more eco-
friendly approach, i.e., a solvent-free method. Starch was thoroughly
dried in a vacuum oven at 70 °C for 48 h. After that, OSA (10 wt % of
the starch mass) was mixed gently with the dried starch, and then the
mixture was kept in the preheated oven at 100 °C for 24 h. The
surface-treated starch was assigned as OSA-g-starch and used for
blending with the copolyamide matrix. The established reaction
between the OSA molecules and starch particles is schematically
depicted in Figure 1b. To characterize the surface-treated particles, we
dispersed OSA-g-starch in toluene at 90 °C and mixed for 1 h. After
that, it was filtered and washed with ethanol (Etax B) to remove any
unreacted OSA residue.
Blending and Injection Molding. Dierent amounts of OSA-g-
starch were melt-blended with the copolyamide in a counter-rotating
twin-screw extruder (Brabender Plasti-Corder PLE 651 with DSK 42/
7 twin-screw extruder, The Netherlands) where the screw speed was
20 rpm, and the heating zone temperatures were set to 155 °C (feed),
160 °C (middle), and 160 °C (Die). The output filament was
solidified under an air atmosphere and then cut by using a pelletizer.
The pellets were then fed to an injection molding machine (Engel ES
200/40) to prepare the tensile test specimens. A gentle dosing
procedure was used to avoid any damage to the starch; thereby, the
four heating zones were set at 155 °C (feed), 160, 165, and 160 °C
(die) while a screw speed of 30% (counter-pressure 1 bar) was used
during the dosing step. The injection speed of 150 mm s−1was used
during the injection step, followed by after pressure (20 s at 15 bar).
The mold temperature was set at 40 °C. For other characterizations,
the pellets were hot-pressed for 2 min using a Fontijne Lab Press-TP
(The Netherlands) at 160 °C and then cold-pressed at 20 °C for 10
min. The portions of the OSA-g-starch in the copolyamide matrix
were selected as 10, 30, 50, and 70 wt %, and the samples were coded
as PSMS10, PSMS30, PSMS50, and PSMS70, respectively. Bio-
composite containing 50 wt % native starch (PNS50) was also
prepared to compare its properties with biocomposites containing
surface-treated starch.
Plain copolyamide and PASMS10 and PASMS50 composites were
chosen for filament production to create actuators. The materials were
introduced into a twin-screw microcompounder (Xplore Instruments
Midi Extruder). The temperature was set at 160 °C, and the rotational
speed was maintained at 15 rpm. The extruded material was solidified
on a conveyor operating at a maximum speed of 10 m/min. It should
be highlighted that even at a high starch content of 50 wt %, no
filament breaking was observed during the extrusion process. The
resulting filament was wound around a spindle and utilized for coiled
actuator fabrication, as detailed in a subsequent section. The filament
diameters for plain copolyamide, PASMS10, and PASMS50 were 190
±5, 270 ±10, and 420 ±20 μm, respectively.
Fabrication of Coiled Actuators. PSMS50 was chosen for
preparing polymer actuators due to the higher shape fixity (Rf) and
shape recovery (Rr), as will be demonstrated in the following section,
compared with the other samples. The actuators were fabricated with
self-made equipment that resembled the fabrication device. First, the
filament twisted until it started to overtwist. The possible maximum
number of twists causes a stronger actuation.
34
After being twisted,
the filament was coiled around a metallic mandrel. By varying the
twisting and coiling directions, it was possible to make contracting
and expanding actuators. An expanding heterochiral structure was
used for testing and comparing the actuator properties because
expansion was more unlimited as a moving direction. Then, the
twisted specimen was thermoset for 150 min at 120 °C in a
hydrothermal oven.
Weaving the Textile Prototype. A 3D fabric design was
developed by the integration of thermoresponsive PSMS50 coiled
actuators into the textile substrate through weaving. The prototype
was a multilayered woven fabric that laid as a flat fabric before
actuation while on thermal actuation, forming a three-dimensional
cellular structure. Figure 2a illustrates the cross-sectional and
schematic representation of the layers, while Figure 2b showcases
the top view of the fabric woven with PSMS50 actuators and cotton in
the first layer. Furthermore, the digital image of the fabricated
prototype as well as the image of fabric photographed in a dark
chamber are presented in Figure 2c,2d, respectively.
The prototype was woven on a Thread Controller 2 (TC2) Digital
Jacquard loom. In a TC2 loom, each warp thread could be
programmed to work independently, allowing for the weaving of
complex, multilayered fabrics and weave patterns. Moreover, while the
warp movements of the loom were computerized, the weft insertion
was done manually by hand, further allowing newly developed
materials, in this case, PSMS50 coiled actuators, to be prototyped
with ease. The five-layered fabric comprised 100% cotton warp and
wefts consisting of Linen-Tencel, Cotton, glow-in-the-dark yarn, and
PSMS50 actuators. The top layer of the fabric was woven with 100%
cotton and PSMS50 actuators, in which the plain weave bonded the
actuators to the textile structure so that they did not escape the woven
layer during actuation. The cotton yarns and the plain weave structure
complemented and aided the actuation of PSMS50 actuators within
the textile network. The cotton yarns provided the required flexibility
to the top layer, allowing the actuators to contract and expand easily.
The weft of the other 4 layers was formed by Linen-Tencel. The linen
weft was conducive to creating the diamond structure, as the rigidity
of the yarn helped to create well-defined cells that retained their shape
without wrinkling during the actuation. This enabled the fabric to
open into a fully 3D form during actuation. The interplay of rigid
yarns and shrinking yarns in weaving has been tested in earlier work
to develop 3D fabrics.
35
The glow stripes in the dark yarns were
intermittently introduced as a supplementary weft in the second and
fourth layers. When the fabric was actuated into its three-dimensional
form, the glow-in-the-dark yarns peeked through to form glowing
interlacing stripes in the dark, as shown in Figure 2c.
Characterization Methods. Fourier transform infrared spectros-
copy (FTIR) was run on a PerkinElmer FTIR with an attenuated total
reflection (ATR) machine to investigate the chemical structure of the
synthesized PA11, PA1218, and copolyamide as well as starch and
OSA-g-starch. The spectra were recorded between 4000 and 500 cm−1
under a scan rate and resolution of 16 and 4 cm−1, respectively. The
proton nuclear magnetic resonance (1H NMR) spectra conducted by
a Bruker AV III 400 NMR spectrometer were used to further evaluate
the chemical structure of the homopolymers and copolymer. A
mixture of chloroform-dand trifluoroacetic anhydride (90/10, V/V)
was used to dissolve the sample before subjection to measurement.
The gel permeation chromatography (GPC) was performed on an
Agilent Multidetector machine to measure the number and average
molecular weights, as well as the polydispersity index of the
homopolymers and copolymer. A mixture of chloroform (for analysis
EMPARTA ACS) and trifluoroacetic anhydride (90/10, V/V) was
used as a solvent. Polystyrene standards, dissolved in the same solvent,
were used for calibration. The degree of OSA substitution was
quantified by an elemental analysis performed on a Thermo Flash
Smart CHNSO Elemental Analyzer. First, the calibration curve was
plotted considering the measured values of carbon, hydrogen, and
nitrogen elements in the native starch. The curve is depicted in Figure
S1, where the measured values were graphed versus the theoretical
values, that is, C as 44.45%, O as 49.34%, and H as 6.21%. The
measured elemental values were then modified due to the calibration
curve. After that, the corrected value of the carbon element was
employed to calculate the degree of substitution (DS), considering eq
1.
C
C
DS 72.06 162.14
350.5 288.26
=
×
×
(1)
where 72.06 is the carbon mass in the anhydroglucose unit, Cis the
carbon concentration in the sample (obtained from the elemental
analysis), 162.14 stands for the molecular weight of the
anhydroglucose unit, and 288.26 and 350.5 are the carbon mass
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D
and OSA molecular weight, respectively. The DS was then used to
calculate the weight percent of grafted OSA concerning the molecular
weight of OSA (350.5 g mol−1) and the anhydroglucose unit (162.14
g mol−1). The reported values were the mean of three replicates ±the
error. Dierent crystallization properties of the homopolymers and
copolymer, including melting point (Tm), crystallization temperature
(Tc), melting enthalpy (ΔHm), crystallization enthalpy (ΔHc), and
degree of crystallinity (χc), were obtained by dierential scanning
calorimetry (DSC) analysis performed on a TA Instruments MT-DSC
Q2000 machine. A two-cycle method was applied under a nitrogen
atmosphere, where the thermal history of the sample was removed
through the first heating−cooling cycle, and the aforementioned
properties were extracted from the second cycle. The temperature
range was between −20 and 250 °C, and the heating/cooling scan
rate was fixed at 10 °C min−1. The crystallinity was calculated by using
eq 2, in which ΔH0is the enthalpy of a 100% crystalline sample.
H
H
C
m
0
=
(2)
For copolyamide, ΔH0was calculated considering the weight
percent of homopolymers, i.e., PA11 and PA1218, based on eq 3,
where 28 and 72% were the weight percents of PA11 and PA1218 in
the copolyamide, respectively. The melting enthalpy of a 100%
crystalline PA11 was considered as 226.4 J g−1,
36
but since there was
no value for a 100% crystalline PA1218 in the literature, the enthalpy
of a 100% crystalline PA1212 (292.2 J g−1) was used.
37
H H H0.28 0.72
copolyamide
0
PA11
0
PA1212
0
= × + ×
(3)
DSC was further used to investigate the crystallization properties of
the developed composites. A similar thermal cycling test was
performed; just the degree of crystallinity was calculated using eq 4,
where xis the weight percent of filler particles, e.g., OSA-g-starch.
H
H x(1 )
C
m
0
=
(4)
The grafting of OSA on starch was qualitatively investigated by
thermogravimetric analysis (TGA) thermograms performed on a TA
Instruments TGA Q500 machine (Figure S2). Furthermore, the
thermal decomposition of the developed biocomposites was
monitored by TGA (Figure S3 and Table S1). The sample was
heated under a nitrogen atmosphere from room temperature to 800
°C at a heating rate of 10 °C min−1. The morphology of the starch
particles before and after OSA grafting, as well as the dispersion level
of OSA-g-starch into the polymer matrix, were evaluated by scanning
Figure 3. (a) FTIR spectra, (b) 1H NMR spectra, and (c) DSC thermograms of the synthesized homopolymers and copolymer. (d) FTIR spectra
and (e, f) SEM images of starch before and after treatment with OSA.
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E
electron microscopy (SEM) images taken by a Zeiss Sigma VP
(Entry-level SEM) machine. The imaging was done from the surface
of the particles, while it was carried out from the cryofracture cross-
sectional area of plain matrix and biocomposites. A thin layer of gold−
palladium was sputtered on the sample surface prior to the subjection
of SEM imaging. The mechanical properties of the samples were
evaluated by tensile testing done by a Universal Tester Instron 5944
machine. Tensile modulus, yield stress, tensile strength, tensile strain
(or elongation at break), and toughness, i.e., the area under the
stress−strain curve, were extracted and reported. The test was
conducted based on ASTM D638−02, with a stretching rate of 5 mm
min−1. The samples were conditioned for 48 h at a temperature of 23
°C and a relative humidity of 50%. Each measurement was repeated 5
times, and the mean value ±standard deviation was reported. The
thermomechanical characteristics of the samples, including the storage
modulus (E′), loss modulus (E″), and loss factor (tan δ), were
evaluated versus temperature by dynamic mechanical analysis (DMA)
conducted on a TA Instruments DMA Q800 machine under an air
atmosphere. The temperature was ramped up from −20 to 140 °C
with a rate of 5 °C min−1. The preload, frequency, and strain rate were
1 N, 5 Hz, and 1%, respectively.
37,38
The glass transition temperature
(Tg) was extracted from the tan δcurve.
The shape memory eect of the plain matrix and biocomposites
was characterized by Q800 DMA using a controlled-force mode. A
strip-shape sample with the dimensions of 5 mm ×3 mm ×0.5 mm (l
×w×t) was heated to 80 °C with a heating rate of 5 °C min−1and
kept 5 min at this temperature. Then, the force was applied to 18 N at
an increasing rate of 0.5 N min−1. Subsequently, the sample was
cooled to −30 °C followed by a 5 min isothermal step at this
temperature. After that, the temporary shape was recovered by
removing the load with the rate of −0.5 N min−1until reaching the
minimum force of 0.1 N. This process was repeated for 4 cycles. Rr
and Rfwere calculated based on eqs 5 and 6, respectively.
39
R N
N N
N N
( ) ( ) ( )
( ) ( 1) 100%
r
m p
m p
= ×
(5)
R N
N
N
( ) ( )
( ) 100%
f
u
m
= ×
(6)
where εm,εp,εu, and Nindicate the strain after stretching (before
cooling), strain after recovery, strain in the fixed temporary shape, and
cycle number, respectively.
39
It should be highlighted that the upper
limit for temperature was set at 80 °C for the plain copolyamide
matrix and PSMS10; however, it was 70 °C for PSMS50 because, at
higher temperatures, the elongation of the sample was higher than the
operating range of the device. The viscosity of the samples and their
viscoelastic performance were investigated by melt rheology testing in
an oscillatory mode on an Anton Paar Physica MCR 301 machine.
The tests were conducted at 160 °C, the same temperature used for
melt blending, with parallel geometry (PP25). A strain sweep test was
performed from 0.01 to 100% at a fixed angular frequency of 1 Hz to
find the linear viscoelastic region. Afterward, the melt strength and
flowability were investigated within the linear viscoelastic region, i.e., a
fixed strain rate of 1%, through a frequency sweep test ranging from
0.01 to 100 Hz. The trend of both moduli, i.e., storage modulus (G′)
and loss modulus (G″), as well as the complex viscosity (|η*|) versus
angular frequency was plotted and explored.
■RESULTS AND DISCUSSION
Copolymerization. FTIR spectroscopy was employed as a
fundamental analytical tool to investigate the chemical
structures of the homopolymers and copolyamide. As shown
in Figure 3a, both homopolymers and copolyamide presented
typical polyamides FTIR characteristic peaks. Namely, the N−
H group provided a broad bond at 3300 cm−1,−CH2
asymmetric stretch and symmetric stretch peaks appeared,
respectively, at 2917 and 2847 cm−1, amide I (CO) and
amide II (−NH−CO−) presented two peaks at 1633 and 1541
cm−1, CO bending formed a bond at 1465 cm−1, and a bond
at 938 cm−1originated from amide IV.
40,41
It is noteworthy to
mention that the N−H bond peak intensity at 3291 cm−1
reduced considerably in PA1218 and copolyamide, indicating a
decrease in the amount of amide functional groups, suggesting
an increase in the aliphatic segments’ chain length.
11,42,43
1H NMR spectra were used to determine the structure of the
synthesized homopolymers and copolyamide concerning the
hydrogen-1 nuclei. The results are plotted in Figure 3b. The
peak signals at around 10.5 and 7.3 ppm were attributed to the
trifluoroacetic anhydride and chloroform-d, respectively. The
other resonances that appeared in the range of 1−4 ppm were
characteristic peak signals for polyamide
37,38,44
proving the
successful synthesis of the PA11 and PA1218 homopolymers
as well as the copolymer through the employed polycondensa-
tion reaction. Specifically, the resonance at approximately 3.67
ppm (marked a in Figure 3b) originated from the −CH2
proton adjacent to the amino groups, the peak signal at
approximately 2.7 ppm (marked b in Figure 3b) was attributed
to the −CH2proton adjacent to the carbonyl groups, and the
peaks in the range of 1.2 to 1.7 ppm (marked a and d in Figure
3b) were assigned to the −CH2proton in the aliphatic chain.
One of the strengths of the synthesized copolyamide is its
relatively low melting point, making it an exciting polymer for
compounding with natural fillers, e.g., starch. As such, the
melting points of the synthesized homopolymer and
copolymer were measured by DSC. The DSC scans are
illustrated in Figure 3c. Furthermore, the relevant DSC data
extracted from the curves are summarized in Table 1. All
polymers showed double melting peaks, a dominant peak
accompanied by a minor adjacent shoulder, indicating a
polymorphic structure. The appearance of double melting
peaks might suggest the melting of dierent crystalline phases,
i.e., α- and γ-forms, or the fusion of the same crystal phase, α-
form, yet with varying thicknesses.
10,45,46
Furthermore, all
samples exhibited an exothermic peak corresponding to their
crystallization temperature, indicating the formation of a single
crystalline phase, even in the copolymer. Noticeably,
copolyamide had its own crystallization temperature, not
those that appeared for its constituents, indicating that the
monomers were not polymerized separately.
47
This was further
supported by GPC results, Table 1 and Figure S4, where the
copolyamide had a unimodal molecular weight distribution
rather than a bimodal one.
Table 1. DSC Data and Dierent Molecular Weights of the Synthesized Homopolymers and Copolyamide
sample Tc(°C) ΔHc(J g−1)Tm(°C) ΔHm(J g−1)χC(%) Mn(g mol−1)Mw(g mol−1) PDI
PA11 163 41.20 183, 190 41.92 18.20 53,700 91,500 1.70
PA1218 144 52.41 165, 171 32.25 11.04 87,900 15,800 1.79
copolyamide 110 36.42 135, 142 29.70 10.85
a
72,000 18,600 2.60
a
The melting enthalpy of 100% crystalline copolyamide was considered 273.77 J g−1due to eq 3.
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On the other hand, copolymer melted and crystallized at
relatively lower temperatures than its parent polymers. In
addition, its melting and crystallization enthalpies were lower
than those of the homopolymers. These results suggest that the
copolymerization of dierent monomers/polymers aected the
chain organization, changed the polymer chains’ regularity, and
inhibited crystallization in the copolymer, which resulted in
lower crystallinity and melting point.
48−50
Starch Surface Modification. FTIR spectra were used to
investigate the possible covalent bonding between starch and
OSA molecules. As plotted in Figure 3d, starch and OSA-g-
starch exhibited some common peaks, which were character-
istic of starch.
26,51
Specifically, a broad band at 3270 cm−1
originated from the intra- and intermolecular hydrogen bonds
of O−H stretching, the peaks between 3000 to 2800 cm−1
resulted from −CH stretching bonds, a peak at 1643 cm−1
assigned to the water molecules trapped in the noncrystalline
region of the sample, the peaks located at 927, 1010, 1079, and
1148 cm−1attributed to the −C−O−C−bonds of the
anhydroglucose unit, and the bands at 574, 769, and 863 cm−1
arose from the C−C stretching and C−H bending vibrations
of the glucosidic ring.
The relative intensities of some peaks changed in the OSA-g-
starch samples. In addition, some new bands appeared,
indicating the grafting of the OSA molecules on starch via a
covalent reaction between the hydroxyl group of starch and
anhydride rings of OSA. For instance, the grafted alkyl chain
gave rise to two new peaks at around 2916 and 2850 cm−1in
the surface-treated starch. Besides, the relative peak intensity
corresponding to hydroxyl groups at around 3300 cm−1
decreased significantly, suggesting a lower concentration of
hydroxyl groups in the surface-treated sample attributed to
their interaction with OSA molecules via esterification.
Notably, a newly formed band at 1740 cm−1, which could be
considered the characteristic of the carbonyl groups in the
ester bonds (CO),
26,38,52
further proved the claim of the
esterification reaction. When the spectra are interpreted, it
should be taken into account that the C−C−O stretching and
O−C−C stretching bonds of the ester groups at 1240 and
1047 cm−1possibly overlapped with the characteristic peaks of
the anhydroglucose unit.
The degree of substitution and the weight percent of the
grafted OSA, were estimated based on the elemental analysis.
The obtained data are summarized in Table 2. The DS was
approximately 5.45 and 2.34 per 100 anhydroglucose units,
corresponding to approximately 9.94 and 4.26 wt % OSA in
the surface-treated starch before and after washing, respec-
tively. The OSA concentration before washing the sample had
Table 2. Elemental Analysis Results of Starch and OSA-g-Starch
sample carbon (%) hydrogen (%) nitrogen (%) oxygen (%) sulfur (%) DS
a
OSA (wt %)
Theoretical Values Regarding the Anhydroglucose Formula, C6H10O5
starch 44.64 6.19 0 49.16 0
Measured Values
Starch 41.33 ±0.34 6.18 ±0.09 0 47.55 ±0.69 0
OSA-g-starch
b
45.94 ±0.81 6.92 ±0.08 0 45.53 ±0.38 0
OSA-g-starch
c
43.89 ±0.09 6.58 ±0.02 0 47.27 ±0.08 0
Corrected Values
d
Starch 39.91 6.32 0 46.79 0
OSA-g-starch
b
48.43 7.10 0 48.00 0 5.45 9.94
OSA-g-starch
c
46.26 6.74 0 49.83 0 2.34 4.26
a
Calculated from eq 1.
b
Before washing.
c
After washing.
d
Regarding the calibration curve (Figure S1).
Figure 4. SEM images from cryofracture surface area of (a) neat copolyamide, (b) PSMS10, (c) PSMS30, (d) PSMS50, (e) PSMS70, and (f)
PNS50 with 500×magnification.
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an excellent agreement with the added value (10%), as
previously explained in the Experimental Section. Approx-
imately half of the added OSA was washed away during the
washing process, indicating that the remaining OSA reacted
covalently with starch molecules, as observed in the FTIR
spectra.
The eect of surface treatment on the morphology of starch
particles was monitored using SEM imaging. The micrographs
are shown in Figure 3e,3f. The starch granules had an oval
geometry with an average diameter of 35 ±15 μm, in line with
the value reported in the literature for dierent starch types.
53
Besides, their surface was smooth. After surface treatment, the
particle size did not change considerably; however, the surface
was no longer smooth and became rough, suggesting a
successful grafting of OSA molecules.
54−56
Biocomposites Morphology. The microstructure of the
biocomposites was evaluated by using cryofracture cross-
sectional SEM imaging. The SEM micrographs of the
biocomposites, as well as the plain copolyamide and the
biocomposite containing 50 wt % of unmodified starch, i.e.,
PNS50, are shown in Figure 4. The unmodified copolyamide
surface was smooth without any cracks and particles, indicating
a plain polymer with a tough continuous surface rather than a
brittle one. Likewise, no significant surface defects, i.e., pores
or cracks, were detected in biocomposites containing surface-
treated starch, while OSA-g-starch distributed uniformly into
the polyamide matrix at a filler content up to 50 wt % with no
sign of particle agglomeration. These observations indicated
good miscibility of copolyamide and surface-treated starch and
improved compatibility and interfacial adhesion between
phases,
18,57−59
achieved via the grafting of OSA on the starch
surface.
In contrast, the PNS50 composite containing 50 wt %
unmodified starch exhibited significant surface defects along
with the formation of voids, confirming that those two
substances were inherently incompatible with poor interfacial
adhesion between the phases.
60,61
As can be observed in Figure
4e, the starch reversed from the dispersed phase to the
continuous one in PSMS70 with 70 wt % of OSA-g-starch,
which counts as another piece of evidence of the excellent
miscibility of the two phases.
18
Overall, the average diameter of the starch granules was
about 30 μm, which is in good agreement with what was
observed previously for the plain OSA-g-starch. This indicates
that the starch particles were not destroyed under the applied
shear stress in the extruder, which has already been introduced
in the literature as another reason for the absence of any cracks
and folds on the cross-sectional area of the biocomposites.
62
It
is worth notifying that some researchers observed significant
agglomerates and defects in the composites with significantly
lower starch loadings,
53,63,64
highlighting the advantage of the
employed method for the surface treatment of starch. In this
study, successfully incorporating such a high loading of starch
particles in the polyamide matrix introduces an eco-friendly,
green, durable, and sustainable plastic that may significantly
reduce the carbon footprint and greenhouse gas emissions.
65,66
Crystallinity Study. The impact of OSA-g-starch particles
on the crystallization behavior of the developed composites
was thoroughly investigated by using DSC. The DSC curves
are illustrated in Figure S5, and the corresponding thermal
transition data are summarized in Table 3. Interestingly, the
composites with up to 30 wt % surface-modified starch
particles exhibited double melting peaks akin to those observed
in the copolymer. This suggests that the crystal structure
remained relatively unchanged at these concentrations.
However, at higher starch contents, only a single melting
peak was detected. This intriguing phenomenon indicates a
transformation from a less ordered γ-crystalline form to a more
highly ordered and densely packed α-crystalline structure.
67
Alternatively, it suggests the formation of the α-crystalline form
with a relatively uniform thickness.
Furthermore, the presence of high concentrations of starch
led to an increase in the melting point. This can be attributed
to the strong hydrogen bonding between starch particles and
polymer chains, which restricted the chain mobility and
delayed the melting of the polymer chains. The crystallization
temperature also shifted to higher values, indicating that the
starch particles acted as nucleation sites during polymer
crystallization from the melt, lowering the energy barrier of
nucleation. As a result, less supercooling of the melt was
required to initiate crystallization, leading to a shift in the
crystallization temperature to higher values.
68
Previous studies
have reported that higher crystallization temperatures can lead
to increased crystallinity in the polymer matrix.
37
Conse-
quently, the degree of crystallinity (χc) increased from 10.85%
in the plain copolymer to 12.96% in the PSMS50 composite.
This enhancement confirms that the granular dispersed
morphology favored the crystallization of copolyamide and
that the nucleation eect was reinforced by the presence of
surface-modified starch particles.
22,69
Moreover, the higher melting point observed in composites
with filler contents higher than those of the plain matrix can
also be attributed to their relatively higher crystallinity or
change in the crystalline phase. It is worth noting that the
relatively higher crystallinity in the composites is advantageous,
as it can lead to a higher stiness and tensile modulus. This is
because crystalline regions provide stronger and more ecient
load transfer pathways.
70
Additionally, crystalline regions play a
significant role in enabling and controlling the shape memory
eect, which could prove beneficial in these materials.
71
These
claims and findings will be further investigated in the
subsequent sections to gain a deeper understanding of the
overall composite behavior and its potential application.
Mechanical Properties. The mechanical properties are
important for a material to be applied, among other things, in
the fabrication of textiles. Therefore, dierent mechanical
characteristics of the synthesized copolyamide and developed
biocomposites, including the tensile modulus, yield stress,
tensile strength, tensile strain, and toughness, were evaluated.
The typical stress−strain curves are depicted in Figure 5a.
Furthermore, the aforementioned properties are compared in
Figure 5b and Table 4. The mechanical properties of the
PNS50 biocomposites are also provided for comparison. The
Table 3. DSC Data of the Synthesized Copolyamide and the
Developed Composites
sample Tc
(°C)
ΔHc
(J g−1)Tm(°C)
ΔHm
(J g−1)
χC
a
(%)
copolyamide 110 36.4 135, 142 29.7 10.9
PSMS10 111 26.8 135, 142 28.3 11.5
PSMS30 115 22.9 135, 143 22.7 11.9
PSMS50 119 18.1 149 17.8 13.0
PSMS70 115 12.4 146 10.0 12.3
a
The melting enthalpy of 100% crystalline copolyamide was
considered 273.77 J g−1due to eq 3.
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plain copolyamide showed a very high elongation of about 420
±20%, with tensile modulus and tensile strength of 499 ±24
and 40.3 ±1.9 MPa, respectively. While the tensile strength of
the copolyamide was in the range reported for commercial
PA11
72,73
and long-chain aliphatic polyamides
37,42,43
the
tensile modulus was significantly lower. In other words, the
copolyamide behaved more like a strong and soft polymer,
74
which could be an advantage for blending with mostly rigid
biobased fillers like starch.
75,76
Accordingly, the yield stress,
tensile strength, tensile strain, and toughness were systemati-
cally reduced in the biocomposites, while the tensile modulus
increased. For instance, in the biocomposite containing 50%
surface-modified starch, i.e., PSMS50, tensile strength and
tensile strain decreased, respectively, by ∼68 and ∼45%,
probably due to the brittleness, rigidity, and poor mechanical
properties of starch compared to polyamide. At the same time,
the tensile modulus was enhanced by approximately 25%, a
notable improvement that can be attributed to the higher
crystallinity observed in the composites compared to that of
the plain matrix. Crystalline regions within a polymer matrix
play a pivotal role in reinforcing mechanical properties,
particularly by increasing the material’s stiness. The align-
ment and close packing of polymer chains in these crystalline
regions facilitate superior load transfer and bolster resistance to
deformation, ultimately culminating in enhanced stiness.
77
It
is worth notifying that the tensile strain dramatically dropped
to 10% in the biocomposite containing 50 wt % native starch,
i.e., PNS50, proposing a brittle composite. In contrast, it was
more than 200% in PSMS70 with an even higher surface-
treated filler loading. The photograph captured from bent
PSMS70 (Figure 5c) clearly shows the flexibility of this sample.
Altogether, the mechanical properties establish the benefit of
the employed surface modification method for making starch
particles compatible with the polyamide matrix.
The trend of copolyamide’s mechanical properties as a
function of starch concentration was compared with the
research works performed on starch-based thermoplastic
composites. A few researchers claimed improvement in the
mechanical properties of the polymer matrix after melt
blending with starch. For instance, Yuso et al.
60
prepared
Figure 5. (a) Typical stress−strain curves, (b) comparison of dierent mechanical properties of the synthesized copolyamide and biocomposites,
(c) photograph of bent PSMS70 biocomposite, (d) storage (E′) and loss (E″) moduli, and (e) loss factor (tan δ) of the synthesized copolyamide
and biocomposites versus temperature. The solid and blank symbols represent the storage modulus and loss modulus, respectively.
Table 4. Mechanical Properties of the Synthesized Copolyamide and the Developed Biocomposites
sample tensile modulus (MPa) yield stress (MPa) tensile strength (MPa) tensile strain (%) toughness (MJ cm−3)
copolyamide 499 ±24 22.7 ±1.1 40.3 ±1.9 420 ±20 6.7 ±0.3
PSMS10 576 ±21 19.3 ±0.7 34.9 ±1.3 440 ±15 5.3 ±0.2
PSMS30 613 ±26 17.8 ±0.7 23.7 ±1.0 320 ±14 3.0 ±0.1
PSMS50 627 ±35 13.1 ±0.8 12.9 ±0.7 230 ±13 1.3 ±0.01
PSMS70 872 ±28 8.6 ±0.2 9.8 ±0.3 205 ±7 0.9 ±0.01
PNS50 654 ±38 10.7 ±0.6 9.5 ±0.5 10 ±0.6 0.05 ±0.01
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PLA/thermoplastic starch with dierent compositions and
reported around 35% improvement in both tensile modulus
and tensile strength of biocomposite with 30% starch,
attributed to the good matrix−filler interaction. Nevertheless,
similar to our findings, most of the previous studies revealed a
significant reduction in the mechanical performance of the
matrix. Namely, Landreau et al.
78
observed an approximately
70% reduction in the tensile strength of the PA11 matrix
containing 50 wt % starch content compatibilized by sodium
carboxymethylcellulose. Furthermore, Noivoil et al.
23
used
oligo(lactic acid)-grafted starch as a compatibilizer and
prepared a PLA/thermoplastic starch biocomposite blend
(50/50) with dierent concentrations of compatibilizer. They
reported a 50% reduction in both tensile strength and tensile
modulus of the biocomposite but a 9% increase in tensile
strain. Similarly, a significant decrease in Young’s modulus and
tensile strength of PLA has been reported after blending the
matrix with thermoplastic starch.
79,80
Likewise, Zhang et al.
81
prepared PLA/starch composites (55/45) compatibilized by
maleic anhydride and observed a 15 and 10% reduction in
tensile strength and tensile strain, even in the biocomposite
containing 2% maleic anhydride as a compatibilizer. Moreover,
Zhou et al.
63
observed a 13% decrease in the tensile strength of
PVA/corn starch (80/20) biocomposite.
Thermomechanical Performance. To further evaluate
the OSA-g-starch dispersion into the copolyamide matrix, the
trends of dynamic storage modulus (E′) and dynamic loss
modulus (E″), as well as the damping factor (tan δ), were
evaluated as a function of temperature from −20 to 140 °C by
dynamic mechanical analysis (DMA). The results are
presented in Figure 5d,5e. Although the moduli were constant
at low temperatures, they experienced a sharp reduction of
about an order of magnitude between 40 to 60 °C, associated
with a peak on tan δcurves, indicating a transition from a
glassy state to a rubbery one, also known as glass transition
temperature (Tg). This sharp reduction of moduli around Tg
can be explained by the fact that at low temperatures, the
molecular segmental motions were frozen and thereby caused a
high level of E′from 900 to 1900 MPa. Upon increasing the
temperature, the rigid segmental structure relaxed gently,
resulting in higher molecular chain motions in the system.
82
On the other hand, below Tg, both E′and E″increased
systematically with the increase in the OSA-g-starch content,
which could be due to the hydrogen bonds formed between
the hydroxy group of starch and amine groups of polyamide. It
should be highlighted that the modulus at high temperatures
showed a slight improvement after the matrix was blended with
surface-modified starch particles, particularly in the PSMS50
composite. This enhancement could be highly advantageous
for shape memory applications. The modulus of a material
refers to its stiness or resistance to deformation. In the case of
shape memory polymers (SMPs), the modulus at high
temperatures plays a critical role in determining the recovery
stress and, consequently, the overall performance of SMPs in
various applications.
83−85
When an SMP is deformed above its
switching temperature (Tgor Tm)in this specific instance,
Tgthe high-temperature modulus governs the force it can
generate during shape recovery. Achieving a balance between
the modulus at high and low temperatures is essential for
ecient shape memory behavior. Moreover, the incorporation
of surface-modified starch particles in the PSMS50 composite
has demonstrated the potential for enhancing the high-
temperature modulus, further contributing to improved
shape memory properties, as will be discussed in the following
sections.
As the temperature increased, the flexible long-chain
structure of grafted OSA molecules reduced the rigidity of
biocomposites.
86
Likewise, the peak in tan δcurves, Tg, shifted
to lower temperatures upon increasing OSA-g-starch content.
Namely, Tgshifted from 54 °C in the copolyamide matrix to
48 °C in PSMA50 and then to 40 °C in PSMS70. This
phenomenon suggests that the biocomposites became less
viscous after the temperature was raised, and more molecular
chain mobility was achieved. In other words, the flexibility of
the matrix was developed due to the uniform dispersion of
filler particles, which acted mainly as plasticizers rather than
reinforcing agents.
87,88
The reduction in Tgupon introducing
starch particles might be an advantage for thermoresponsive
shape memory polymer actuators because the actuators could
start to activate at lower temperatures.
A single Tgappears in a fully miscible polymer composite
system, lying between the Tgof the individual polymers.
58,89
The Tgfor potato starch with 13−15% moisture content is
reported to be between 75 and 95 °C.
90
Accordingly, a well-
defined peak in the loss factor indicated eective miscibility
between polyamide and starch as well as a homogeneous blend
microstructure. On the other hand, the height of the tan δ
curve pertains to the matrix/filler interphase internal energy
Figure 6. (a) Storage (G′) and lost (G″) moduli and (b) complex viscosity (|η*|) of the synthesized copolyamide and biocomposites versus angular
frequency at a fixed strain rate of 1% and temperature of 160 °C. The solid and blanked symbols represent storage modulus and loss modulus,
respectively. (c) Digital photograph of the tensile testing specimens extruded at 160 °C (copolyamide and PSMS50) and 220 °C (PNS50).
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dissipation.
87
The height of tan δwas relatively lower in
biocomposites than in the plain matrix, indicating that the
addition of OSA-g-starch particles considerably reduced the
viscoelastic damping factor of the polyamide matrix, and
accordingly, the sample became less rigid.
Rheology and Processability. Rheology measurements
were performed to assess further the microstructural eects of
OSA-g-starch in the copolyamide matrix and gain more insight
into the processing characteristics of the developed bio-
composites. First, the linear viscoelastic region was found
through a strain sweep test conducted at 160 °C and a fixed
angular frequency of 1 Hz. The data are plotted in Figure S6,
where both plain matrix and biocomposites, except PSMS70,
behaved independently of shear strain within the 0.01 to 100%
strain range, indicating a linear viscoelastic region. Accordingly,
a shear strain of 1% was selected for the frequency sweep test
to guarantee that the measurement was in the linear
viscoelastic regime. The frequency sweep results are illustrated
in Figure 6a,6b, where the trend of storage (G′) and loss (G″)
moduli, as well as complex viscosity (|η*|) were depicted as a
function of angular frequency between 0.01 to 100 Hz at 160
°C. G″dominated G′at the test conditions, particularly at
lower frequencies, indicating a fully relaxed state of the
polymer chains, in which the molten polymers revealed more
liquid-like or viscous behavior than solid-like or elastic ones,
91
suggesting good processability of the compounds under the
applied shear stress in the extruder.
The other parameter that reflects the processing flow
performance of composite materials is complex viscosity.
92
As shown in Figure 6b, a Newtonian plateau was observed in
the low frequencies for all of the molten polymers, followed by
a significant reduction at higher applied angular frequencies,
specifying a non-Newtonian behavior known as shear-thinning
viscosity, which could be due to the disentanglement of the
molecular chains.
89,93
This behavior suggests the excellent
fluidity of the molten polymer in the extruder.
94
Many researchers
11,38,89,91,92,95,96
have reported a systematic
increase in both complex viscosity and storage modulus of the
matrix upon adding the filler due to the generation of a
network-like structure inside the molten polymer, which led to
the increase of the molecular entanglements of polymer chains,
as well as owing to the interfacial adhesion and hydrogen
bonding between the filler and polymer chains. Nevertheless,
in this study, the complex viscosity and moduli decreased upon
increasing OSA-g-starch content, indicating the lubricating
eect of surface-treated starch. This trend could be attributed
to the reduced entanglement of polyamide chains obtained by
voluminous OSA molecules, which brought in more free
volume into the polyamide matrix, as well as to the orientation
of starch particles with flexible OSA chains along with the
applied shear forces that promoted the chain mobilities,
behaving thus similarly to previously reported for composites
containing surface-modified lignin
93
and clay
97
particles.
To support the benefit of the surface treatment on the
processability of the polyamide matrix, the viscosity of the
composite containing 50 wt % native starch, i.e., PNS50, was
included in Figure 6b. Contrastingly to modified starch, the
addition of native starch increased the viscosity to at least 1
order of magnitude when compared to the composite with 50
wt % surface-modified starch, i.e., PSMS50. The increased
viscosity made the PNS50 composite more challenging to
process, and we applied a higher temperature (220 °C) to
extrude this sample. As a result, this composite revealed
significant discoloring compared to other biocomposites
processed at 160 °C (Figure 6c), probably due to the partial
decomposition of starch particles at such a high processing
temperature. This observation firmly supports the benefit of
the synthesized low-melting-point copolyamide for blending
with biobased fillers, e.g., starch, as well as the advantage of the
applied surface modification method in making starch particles
compatible with polyamide matrix.
Shape Memory Eect and Thermally Responsive
Fabric Prototype. Shape memory polymers (SMPs) are a
class of smart materials for which a temporary shape can be set,
and the return to a permanent or partially recovered shape is
further driven by exposure to an external stimulus, such as
temperature change, light, voltage, and humidity.
98,99
One-way
shape memory polymers (1W-SMPs), as the name suggests,
have the ability to return to their original or permanent shape
from a fixed temporary shape in a one-time manner. Two-way
shape memory polymers (2W-SMP), on the other hand, can
actuate between two dierent shapes, typically a higher-
temperature (’partially recovered’) shape and a low-temper-
ature shape (temporary shape), upon exposure to the
appropriate stimuli. The partially recovered shape is
determined by geometry-defining netpoints in the polymer
structure, which in the case of polyamide, is expected to be its
crystalline domains.
100,101
Similarly to 1W-SMPs, the polymer
is initially deformed into a temporary shape at a temperature
above its thermal resetting temperature, i.e., at the temperature
that is above the melting point for geometry-defining domains.
The polymer is then cooled under tension to lock it into its
temporary shape, as in 1W-SMPs. To revert to its partially
recovered shape, the 2W-SMP is heated to a transition
temperature significantly below its thermosetting temperature,
and cycling the temperature between low and transition
temperatures will enable repeatable actuation.
99,102,103
High activation temperature and low recovery are major
drawbacks of many thermoplastic SMPs. For instance, Koerner
et al.
104
developed amorphous fluorinated polyimide with a
high shape recovery temperature of 220 °C. Furthermore, Shi
et al.
83
synthesized Tg-based SMP by introducing ionic
moieties to polyether ether ketone (PEEK) with a switching
temperature of around 200 °C. This work was expanded into a
Tm-based SMP by integrating sodium oleate, leading to a
higher switching temperature of 230−240 °C resulting from
the melting of sodium oleate. These limit their utilization in
some applications, such as biomedicine, where fast recovery
and low activation temperature close to the human body are
essential.
101
Additionally, from the point of view of shape
memory textiles, many common textile materials cannot resist
heating to the activation temperatures of the above-mentioned
SMPs. Introducing plasticizers to the polymer matrix can aect
the thermal properties of the SMPs, leading to a decrease in Tg.
For instance, plasticizers can weaken the H-bonds and reduce
the transition temperature. Subsequently, the polymer chain
movement at a specific temperature can increase.
105,106
To this
goal, in the current study, we developed a novel copolyoamide
that presented a relatively low Ttrans at around its Tg, i.e., 55 °C.
Furthermore, we incorporated surface-modified starch particles
into the synthesized copolyamide not only to improve the
green content of the synthesized polyamide but also to reduce
the Tg, which might lead to a decrease in the switching
temperature of the polyamide.
First, the 1W shape memory and shape recovery of the plain
polymer and biocomposites, including 10 and 50% OSA-g-
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starch, were tested via a cyclic thermomechanical test. The
recovery temperature was chosen to be Tg+ 10 °C to allow for
as complete recovery as possible. The results are provided in
Figure 7 and Table S2. In the pure copolyamide matrix, Rrand
Rfwere, respectively, 85.13 ±4.99 and 89.39 ±2.68%,
indicating excellent shape recovery properties of the newly
synthesized copolyamide. The calculated values for both Rrand
Rfwere as high as reported for common shape memory
polymers.
107,108
High storage modulus below Tgand excellent
rubber elasticity above Tgcontributed to high Rrand Rf. After 4
cycles, the film could still return to its original shape, indicating
an excellent thermally induced shape memory.
109
The eect of processing history could be seen in the first
cycle; however, both Rrand Rfimproved by increasing the
number of test cycles, leading to the better shape memory
eect resulting from memory stress reduction and shape
memory training.
39,46
During the first cycle, thermal shape
recovery leads to a reduction in the entropy and disentangle-
ment of some entanglements, diminishing the ability of
polymer chains to return to their previous random coils.
This stretching and loosening of physical entanglements result
in fewer remaining entanglements available for disentangle-
ment in subsequent cycles of thermomechanical loading.
Consequently, the shape recovery value is lower during the first
cycle compared to subsequent cycles.
110,111
Both biocomposites, i.e., PSMS10 and PSMS50, revealed
qualitatively similar behavior over heating/cooling cycles with
relatively high 1W Rrand Rfvalues comparable with those of
the unmodified copolyamide matrix. This confirms that the
presence of starch particles did not interrupt polyamide chains’
mobility even at a relatively high filler content of 50 wt %,
chiefly thanks to the employed compatibilization process. Both
Rrand Rfwere slightly enhanced in the biocomposites,
demonstrating that the OSA-g-starch particles enhanced the
mobility of the polymer chains and helped them easily arrange
and recover to their original shape by disrupting the interchain
supramolecular interactions. The degree of crystallinity (χc) in
shape memory polymers significantly influences their shape
memory properties. Increased crystallinity content often results
in a more pronounced and dependable shape memory
eect.
112
Consequently, the PSMS50 composite with relatively
higher crystallinity exhibits slightly higher Rrand Rfvalues.
However, it is essential to avoid excessive crystallinity, as it
may render the polymer overly rigid, thus limiting its
deformability and shape recovery capabilities. It is noteworthy
that the shape recovery results for PSMS50 were obtained at a
lower operational temperature, i.e., 70 °C. In other words, the
incorporation of surface-modified starch particles not only
improved the shape memory performance of the polyamide
matrix but also significantly reduced the transition temper-
ature, making the composite interesting for low-temperature
actuating applications.
Owing to these biocomposites exhibited two distinct thermal
transition temperatures, Tgand Tm, the 2W shape memory
eect was investigated. To showcase the potential application
of the developed biocomposite in textiles, a heat-responsive
fabric was created using thermos-responsive 2W shape
memory PSMS50 actuators. The fabrication process involved
the use of a custom-built twister-coiler device and thermoset-
ting at a higher temperature of 120 °C, enabling the
demonstration of repeatable thermal actuation. The 2W SM
actuation of PSMS50 coils and the integrated heat-responsive
fabric are depicted in Video S1,Video S2,Video S3, and Video
S4. The choice of using coiled actuators is based on amplifying
Figure 7. Strain/temperature versus time corresponds to the 4 shape memory cycles for (a) copolyamide, (b) PSMS10, and (c) PSMS50. (d)
Shape recovery (Rr) and shape fixity (Rf) percentage for the plain copolyamide matrix and PSMS10 and PSMS50 biocomposites.
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ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
L
the thermally induced torsional motion of twisted filament
when the filament is coiled, as explained previously by Haines
et al.
5
Additionally, Figure 8 illustrates their response to
heating at specific time points.
Upon heating, the shape-morphing fabric exhibited rapid
and significant heat-induced deformation, followed by
successful recovery to its initial dimensions during the cooling
process. Infrared thermal camera recordings, along with Video
S3 and Video S4, showed that the shape-changing deformation
initiated at approximately 50 °C, with the greatest range of
movement achieved at around 100 °C. Notably, the fabric
endured 5 heating/cooling cycles without any change in the
actuation/recovery performance, confirming the durability of
this smart textile.
The cotton yarns incorporated into the fabric provided the
necessary flexibility to the top layer, facilitating contraction and
expansion of the actuators. In essence, the cotton yarns, along
with the plain weave structure, synergistically complemented
and facilitated the actuation of the PSMS50 actuators within
the textile network. This design allowed for the heat-driven
alteration of a 3D textile and could be further utilized to reveal
other hidden functionalities of functional yarns integrated into
the nonactuated textile structure. Consequently, this con-
struction concept enables the gradual exposure of new textile
functionalities based on temperature variations.
In summary, the integration of the developed shape memory
polymer actuators from the PSMS50 biocomposite into the
textile structure showcased significant and heat-induced
deformation. This approach provides a straightforward path-
way for designing dynamic shape-changing textiles that
enhance the capabilities of the current smart textiles.
■CONCLUSIONS
This study highlights the potential of utilizing biobased
materials to address sustainability concerns associated with
petroleum-based fibers in functional textile applications. A
series of biocomposites consisting of surface-modified starch
and polyamide were successfully developed. To prevent the
thermal degradation of starch particles during the melt
blending process, a novel low-melting-point copolyamide was
synthesized by incorporating monomers with dierent alkyl
chain lengths through copolymerization. The resulting
copolyamide exhibited a significantly low melting point of
135 °C. To enhance the compatibility between starch particles
and the copolyamide matrix, a simple, solvent-free method
involving the grafting of octenyl succinic anhydride (OSA)
onto the starch surface through esterification was employed.
The successful grafting of the OSA was confirmed through
various analyses, including FTIR, TGA, elemental analysis, and
SEM images. Subsequently, dierent concentrations of OSA-g-
starch, including a high concentration of 70 wt %, were
successfully melt-blended with the copolyamide matrix at a
relatively low processing temperature of 160 °C. The
developed biocomposites were thoroughly characterized, and
it was observed that the OSA-g-starch dispersed uniformly
within the polymer matrix, resulting in improved interfacial
adhesion and good compatibility. The biocomposites exhibited
excellent stiness/toughness balance, along with remarkable
rheological properties and processability, even at a high filler
loading of 70 wt %. Moreover, the shape memory and shape
recovery properties of the pure polymer and biocomposites,
including those with 10 and 50% OSA-g-starch, were evaluated
through cyclic thermomechanical testing, confirming their
outstanding shape recovery capabilities. Importantly, it was
demonstrated that the shape memory behavior of the
polyamide could be tailored and the actuation temperature
could be reduced by increasing the starch particle content. In
conclusion, this work presents the concept of developing
responsive fabrics by integrating thermoresponsive shape
memory polymers into textiles, thereby adding value to
traditional textiles available in the market and paving the way
for the advancement of smart and functional textiles.
■ASSOCIATED CONTENT
Data Availability Statement
The data that support the findings of this study are available on
request from the corresponding authors. The data are not
publicly available due to privacy or ethical restrictions.
*
sı Supporting Information
The Supporting Information is available free of charge at
https://pubs.acs.org/doi/10.1021/acsami.3c08774.
Starch elements calibration curve was obtained from
elemental analysis results; TGA/derivative TG (TGA/
DTG) thermograms of the starch and OSA-g-starch;
GPC curves of the synthesized homopolymers and
Figure 8. Actuation of heat-responsive fabric using PSMS50 yarn actuators exposed to an IR lamp. (a) IR and (b) digital images of PSMS50 yarn
actuators. (c) IR and (d) digital images of heat-responsive fabric. The fabric undergoes five cycles of heating and cooling. The entire expansion and
contraction process is provided in Video S1,Video S2,Video S3, and Video S4.
ACS Applied Materials & Interfaces www.acsami.org Research Article
https://doi.org/10.1021/acsami.3c08774
ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
M
copolymer; DSC thermograms of the plain copolyamide
and the composites with dierent concentrations of
OSA-g-starch particles; strain sweep test was performed
at a fixed angular frequency of 1 Hz at 160 °C; residue of
the samples at 700 °C, extracted from TGA thermo-
grams; and shape recovery (Rr) and shape fixity (Rf) of
the samples at the second and third cycles (PDF)
Actuation of coils exposed to the infrared heat
(recording was done by an FLIR SC660 thermal
camera) (speed 25×) (MP4)
Actuation of coils exposed to the infrared heat (speed
25×) (MP4)
Actuation of the heat-responsive fabric exposed to the
infrared heat (recording was done by an FLIR SC660
thermal camera) (speed 6×) (MP4)
Actuation of the heat-responsive fabric exposed to the
infrared heat (speed 7×) (MP4)
■AUTHOR INFORMATION
Corresponding Authors
Jaana Vapaavuori −Department of Chemistry and Materials
Science, School of Chemical Engineering, Aalto University,
02150 Espoo, Finland; orcid.org/0000-0002-5923-0789;
Email: jaana.vapaavuori@aalto.fi
Jukka V. Seppälä −Polymer Technology, School of Chemical
Engineering, Aalto University, 02150 Espoo, Finland;
orcid.org/0000-0001-7943-3121; Email: jukka.seppala@
aalto.fi
Authors
Hossein Baniasadi −Polymer Technology, School of Chemical
Engineering, Aalto University, 02150 Espoo, Finland;
orcid.org/0000-0002-0463-337X
Zahra Madani −Department of Chemistry and Materials
Science, School of Chemical Engineering, Aalto University,
02150 Espoo, Finland
Mithila Mohan −Department of Chemistry and Materials
Science, School of Chemical Engineering, Aalto University,
02150 Espoo, Finland
Maija Vaara −Department of Chemistry and Materials
Science, School of Chemical Engineering, Aalto University,
02150 Espoo, Finland
Sami Lipponen −Polymer Technology, School of Chemical
Engineering, Aalto University, 02150 Espoo, Finland
Complete contact information is available at:
https://pubs.acs.org/10.1021/acsami.3c08774
Author Contributions
§
H.B. and Z.M. contributed equally to this work.
Notes
The authors declare no competing financial interest.
■ACKNOWLEDGMENTS
The authors acknowledge the “Academy of Finland” funding
no. 327248 (ValueBiomat) and no. 327865 (Bioeconomy), as
well as funding from NordForsk in the form of the “Beyond
eTextiles” project and from the European research council
project “Autonomously adapting and communicating modular
textiles” no. 949648. The authors also thank Ali Tavakoli for
his eort in editing the videos.
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