Content uploaded by Richard T. Olsson
Author content
All content in this area was uploaded by Richard T. Olsson on Dec 30, 2018
Content may be subject to copyright.
Micromechanics of ultra-toughened
electrospun PMMA/PEO fibres as
revealed by in-situ tensile testing in an
electron microscope
Richard L. Andersson
1
, Valter Stro
¨m
2
, Ulf W. Gedde
1
, Peter E. Mallon
3
, Mikael S. Hedenqvist
1
& Richard T. Olsson
1
1
KTH Royal Institute of Technology, School of Chemical Science and Engineering, Fibre and Polymer Technology, SE-100 44 Stockholm,
Sweden,
2
KTH Royal Institute of Technology, Department of Materials Science and Engineering, SE-100 44 Stockholm, Sweden,
3
Department of Chemistry and Polymer Science, University of Stellenbosch, Private Bag X1, Matieland ZA-7602, South Africa.
A missing cornerstone in the development of tough micro/nano fibre systems is an understanding of the
fibre failure mechanisms, which stems from the limitation in observing the fracture of objects with
dimensions one hundredth of the width of a hair strand. Tensile testing in the electron microscope is herein
adopted to reveal the fracture behaviour of a novel type of toughened electrospun poly(methyl
methacrylate)/poly(ethylene oxide) fibre mats for biomedical applications. These fibres showed a toughness
more than two orders of magnitude greater than that of pristine PMMA fibres. The
in-situ
microscopy
revealed that the toughness were not only dependent on the initial molecular alignment after spinning, but
also on the polymer formulation that could promote further molecular orientation during the formation of
micro/nano-necking. The true fibre strength was greater than 150 MPa, which was considerably higher than
that of the unmodified PMMA (17 MPa). This necking phenomenon was prohibited by high aspect ratio
cellulose nanocrystal fillers in the ultra–tough fibres, leading to a decrease in toughness by more than one
order of magnitude. The reported necking mechanism may have broad implications also within more
traditional melt–spinning research.
Significant efforts are being made to introduce electrospun fibres on a commercial scale using large-scale
electrospinning processing equipment
1,2
. Electrospun ultrathin fibres are currently used in filters and
membranes, and they are also being aimed towards emerging applications such as transparent composites,
scaffold materials for tissue, and templates for cell growth and reinforcement in biomedical silicone materials
1,3–5
.
A major challenge in reaching the targeted applications is the lack of fibre toughness, due to the fine fibre
dimensions that make electrospun fibres sensitive to handle
6,7
. Electrospun fibres have, therefore, mostly been
applied as non-woven mats where the fibres support one another, or collected on a support material as in filter
applications
1
. Attempts to reinforce electrospun polymer fibres by embedding ‘‘rod–like’’ nanofillers which serve
as load-carriers have been reported. Carbon nanotubes (CNTs) have received considerable attention due to their
high strength and high elastic modulus combined with a high aspect ratio .100, but extracted cellulose nanofi-
brils that show more flexibility have also been considered
2,8–12
. Composite fibres with a higher modulus and
strength were obtained in both cases, but the fibres were more brittle
2,10,11
. The absence of toughness was due to the
problematic embedding of the often bent, coiled or spiral-shaped CNTs (occasionally protruding from the surface
of the fibres) and to inadequate filler/matrix adhesion and poor dispersion of the nanofiller inside the polymer
matrix
2,10,11
. Recently, Papkov et al.
6
and Asran et al.
13
demonstrated that tough nanofibres can be obtained by
optimizing molecular orientation/alignment instead of preparing composite fibres. When the diameter of elec-
trospun fibres is reduced (larger draw ratios), an increase in strength and strain-at-break is generally
observed
6,14–21
. However, no direct observation of the fibre failure mechanism has so far been reported for
toughness-improved fibres and their electrospun fibre mats.
In this paper, data for electrospun fibres based on blends of poly(methyl methacrylate) (PMMA) and poly
(ethylene oxide) (PEO), with a toughness more than two orders of magnitude greater than that of pristine PMMA
are presented. Tensile testing was carried out in-situ in the scanning electron microscope on samples consisting of
OPEN
SUBJECT AREAS:
STRUCTURAL PROPERTIES
BIOMEDICAL MATERIALS
NANOCOMPOSITES
POLYMERS
Received
2 June 2014
Accepted
20 August 2014
Published
11 September 2014
Correspondence and
requests for materials
should be addressed to
R.T.O. (rols@kth.se)
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 1
ca. 100 000 parallel fibres. A template transfer method was used to
ensure the non-destructive and reproducible transfer of the fibres to
the micromechanical stage, without any premature stretching of the
fibres
22
after collection of the mats on a rotating cylinder operated at
2000 rpm. The testing revealed that the toughness improvement
depend not only on spinning thin fibres but also on finding a polymer
formulation that promoted extensive fibre necking prior to fracture,
i.e. molecular alignment during fibre stretching. Flexible cellulose
nanofibrils were evaluated as rod-like load-bearers inside the tough-
est PMMA/PEO blend to demonstrate their influence on the fracture
and mechanical fibre performance. The materials presented are bio-
compatible
23–25
, and they offer the possibility to achieve toughened
PMMA-based electrospun fibre systems in applications where non-
harmful formulations are required.
Results and Discussion
Assessment of tensile data for electrospun PMMA/PEO fibre blends.
Fig. 1a shows the stress–strain behaviour of aligned PMMA/PEO
fibres with PEO contents ranging from 0 to 25 wt% (note the split
logarithmic strain axis). In this concentration range the polymers are
miscible
26–29
. All the fibres had a diameter of ca. 2 mm, which allowed
an unbiased analysis of the fibre failure mechanisms of the different
polymer blends. Fig. 1b and Table 1 show that the toughness of the
fibres based on a 5 wt% PEO blend was twice that of the fibres based
on pristine PMMA. A lower concentration of PEO (2 wt%) had
essentially no effect on the tensile properties of the fibres (Table 1).
However, the toughness increased strongly with the content of PEO:
26 and 134 times that of pristine PMMA for 12.5 and 25 wt%,
respectively, and the fracture energy increased from 0.13 MJ/m
3
for
PMMA to 17.1 MJ/m
3
for the blend with 25 wt% PEO. The glass
transition temperature (T
g
) decreased from 124uC for the pristine
PMMA to 75uC for the PMMA blend with 25 wt% PEO,
confirming that a more ductile but glassy fibre material was
formed. It should be noted that both the Young’s modulus and the
tensile strength of the fibres containing 5 and 12.5 wt% PEO were
higher than those of the pristine PMMA fibres. The increase in
Young’s modulus was 18 and 11% and the tensile strength 11 and
77%, respectively. For comparison, the mechanical properties of
solvent-cast films with 25 wt% PEO were evaluated under the same
testing conditions to determine whether the improvement in
toughness was entirely related to the polymer formulation. These
solvent-cast films showed almost the same strength and modulus as
the electrospun fibres (Table 1), but with 61% lower toughness than
the fibres. It was assumed that the greater toughness of the fibres was
solely due to the greater molecular alignment within the fibres,
resulting from the rapid fibre stretching in the electrical field, as
suggested by Papkov et al.
6
However, the in-situ microscopy
characterization revealed that the molecular alignment from the
electrospinning was only part of the explanation of the observed
increase in fibre toughness.
Necking phenomena in toughness-improved PMMA fibres.Fig. 2a
shows the tensile testing curve recorded during in-situ SEM analysis
of the fibre mat consisting of PMMA fibres with 25 wt% PEO. The
testing stage was stopped at different strain levels; 4, 8, 20, 40 and
100% – to investigate the failure behaviour of the fibre mat (Figs. 2 c–
f). A progressive thinning and sequential failure of individual fibres
explained the gradual decrease in stress with increasing strain, when
individual fibres fractured. Higher magnification micrographs of
individual fibres showed that all the fibres deformed via extensive
necking and plastic deformation before failure. Necking was
observed already at the early stage of the deformation (8% strain,
Fig. 2d) and when the strain was increased the necked regions
propagated uniformly along the fibres until complete fibre mat
failure occurred between 130 and 150% strain. The formation of
the necks resulted in an average reduction in fibre cross–sectional
area from 3.8 mm
2
to 0.5 mm
2
(10 observed fibre necks). The true
stress in the fibres necks could therefore be calculated to be ca.
150 MPa before failures started to appear at 40% strain and an
engineering stress of 14.8 MPa.
Fig. 2b shows a close-up inspection of an isolated unloaded fibre at
the early stage of necking, where a 4 mm neck had developed at 20%
strain. The wrinkled surface in the neck was not observed in-situ, and
it is assumed that it originates from an elastic recovery when the fibre
was unloaded for high-resolution microscopy. For comparison,
Fig. 2d displays several 10–15 mm-long necked regions at 8% strain
under load (in-situ). The smaller neck in the unloaded fibre exposed
to an apparently larger strain may originate from an uneven loading
Figure 1
|
(a) Stress-strain curves of fibre mats of PMMA/PEO blends with 0, 5, 12.5 and 25 wt% of PEO (note the split strain-axis to reveal the strain at
break for the fibre mats with the highest toughness. (b) Toughness as a function of PEO content and cellulose content (e: 4 wt% BC, g: 10 wt% BC). (c)
Fibre diameter distribution data of fibre mat samples. (d) Scanning electron micrographs of aligned fibre mats.
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 2
of this particular fibre due to differences in the fibre alignment
(Table 1) or stretching. This highlights the challenge of small-scale
measurements. However, this necking length difference over the
entire 5 mm sample corresponded to a variation of only ca. 0.2%
over the entire mat deformation. In contrast, the necking was always
absent in pristine PMMA fibres because fibre tolerated a strain of
only 1.5%, which resulted in crisp and undeformed fracture surfaces
(Fig. 3b) at a true stress close to that of the entire mat, i.e. ca. 17 MPa.
The shape of the stress-strain curve in Fig. 2a (compared to that in
Fig. 1a) shows stress relaxation, because the tensile stage was stopped
for long periods (total ca. 30 min) to take the micrographs. This
relaxation is common for polymers if the deformation is interrupted
and was not a result of exposure to the electron beam since only a
fraction of the entire mat was exposed. When the load was reinstated,
a slightly higher stress was noted, probably due different relaxations
among the fibres (some fibres were not optimally aligned), which in
turn resulted initially in a more even load distribution over the fibres.
The relaxation was not due to moisture since there was no moisture
in the SEM chamber and a water uptake less than 0.1 wt% was
measured when the samples were removed from the vacuum cham-
ber and placed in an environment with a humidity of 50% RH.
In order to see whether the observed necking phenomenon was a
unique feature of the electrospun fibres, a blend based on the same
PMMA with 25 wt% PEO was extruded into 2.7 mm thick continu-
ous bars for comparison. A mini twin–screw extruder with a short
sleeve die (2 mm) was used to induce minimal molecular orientation
along the bars. The melt-extruded bars showed no signs of necking
during tensile testing at the same strain rate as that used for the
electrospun fibre mats. Only a small narrowing just prior to fracture
was observed and the necking phenomenon was thus unique for the
electrospun fibres. The maximal elongation of the bars was ca. 36%,
with an insignificant toughness increase compared to that of the
same formulation used in the electrospun fibres; see mechanical data
in Table 1 and the photograph of the bar in supplementary informa-
tion, figure S4.
Effect of rod–like cellulose nanofibrils as reinforcement in neck-
ing fibres.In order to investigate the effect of rod-like reinforce-
ments in the electrospun fibres (PMMA with 25 wt% PEO),
bacterial cellulose (BC) nanofibrils were embedded in the fibres at
4 and 10 wt%. The cellulose nanofibrils were extracted from bacterial
cellulose (procedure found in supplementary information S1) and
had a crystal content of 72%
30
. The reported modulus of such
cellulose fibrils/crystals is ca. 100 GPa and they have a strength of
the order of 7–8 GPa
30–35
. The average length of the cellulose fibrils
was 1100 6700 nm, with an aspect ratio (L/h) of ca. 94
30
. The BC
fibrils were dispersed and aligned parallel with the fibre axis in the
electrospun fibres due to their favourable compatibility with the
PMMA, and their flexibility in the oscillating electrospinning jet.
The darker regions in Fig. 3e show the aligned bacterial cellulose
crystals. The composite fibres with 10 wt% BC showed an increase
in the Young’s modulus and tensile strength of 171% and 55%,
Table 1
|
Mechanical data and fibre morphology of electrospun PMMA/PEO fibre mats with different amounts of PEO, data for solvent-cast
films and extruded bars of the PMMA with 25 wt% PEO. Highest values are bold
Material Properties
Fibres
PMMA
100 [%]
Fibres
PMMA/PEO
98/2 [%]
Fibres
PMMA/PEO
95/5 [%]
Fibres
PMMA/PEO
87.5/12.5 [%]
Fibres
PMMA/PEO
75/25 [%]
Films
PMMA/PEO
75/25 [%]
Extruded
PMMA/PEO
75/25 [%]
Toughness [MJ/m
3
] 0.13 60.04 0.13 60.02 0.26 60.02 3.46 60.19 17.1 61.4 6.67 62.0 3.08 60.22
Strength [MPa] 16.8 60.4 16.8 61.3 18.7 62.0 29.8 61.2 18.7 60.4 18.8 60.1 13.9 61.3
Modulus [GPa] 1.49 60.29 1.54 60.03 1.81 60.14 1.65 60.12 0.72 60.05 0.75 60.03 0.69 60.09
Thickness [mm] 1.74 60.35 1.89 60.22 2.54 60.23 1.71 60.25 1.96 60.32 54 63 2.7 60.1 mm
Align. std. dev. 14.3u13.1u14.5u11.6u12.5uN/A N/A
Rel. tough. 1.0 (ref.) 1.0 2.0 27 135 53 24
Figure 2
|
(a) In-situ SEM tensile testing curve of electrospun PMMA with 25 wt% PEO fibres; (b) high resolution micrograph of an early developed fibre
neck; (c) the fibre mat prior to tensile testing; (d) multiple parallel necking of individual fibres at 8% strain; (e) progressive thinning of fibre mat and; (f)
breakage/collapse of aligned and stretched fibres.
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 3
respectively (Table 2 and Fig. 3a (curves 2 and 4)) compared to the
same fibres without cellulose. The electrospun fibres containing 4
wt% BC showed no change in strength but a 25% increase in the
Young’s modulus (Table 2 and Fig. 3a (curves 2 and 3)). However, the
increase in modulus and strength came at the expense of a significant
loss in toughness, by more than one order of magnitude (from 17 to
0.6 MJ/m
3
), Fig. 1b (Table 2). Figure 3a and Table 2 also show the
previously reported increase in strength and modulus of PMMA with
4 wt% nanofibrillated cellulose extracted from wood (WC) (curves 1
and 5)
22
. Electron micrographs of the fractured BC containing fibres
revealed uneven fibre necking and surface porosity development
around the fracture surfaces, and also pull-outs of the cellulose
nanofibrils (Fig. 3d compared to Fig. 3c). Accordingly, even if a
relatively well-dispersed cellulose nanofibre phase was present
(without protruding nanofibres), it appeared that the considerably
stiffer crystals prohibited necking of the electrospun fibres. The
cellulose nanofibrils from wood, with a smaller aspect ratio and a
lower crystallinity did, however, lead to a higher toughness than that
of the brittle pristine PMMA polymer.
Fibre characterization by X-ray diffraction and IR-spectroscopy.
X-ray diffraction was performed to see whether PEO crystals could
have contributed to the improved mechanical properties of the fibres
(or affected the necking characteristics), and IR-dichroism mea-
surements was used to assess chain orientation. It has previously
been shown that PMMA/PEO blends with PEO contents greater
than 25 wt% are biphasic
26
and that PEO crystallizes into PEO-rich
spherulites with PMMA molecules confined in the inter-lamellar
regions
26,36
. Fig. 4a shows diffractograms for the different blends. A
fully X-ray amorphous structure was apparent for concentrations up
to 12.5 wt% PEO, whereas a low crystallinity of 1.2 vol% was detected
in the fibres based on the 25 wt% PEO blend. The samples with
higher amounts of PEO, especially 25 wt% shows a slight shift of
the amorphous halo towards smaller lattice spacing’s, which may be
explained by the lack of bulky side groups in PEO (allowing for
smaller lattice space). Another possibility for this would be the
presence of sub 3 nm crystals appearing amorphous (albeit having
similar lattice spacing as the crystalline phase), which is below the
detection limit for X-ray diffraction
37
.
Electrospun fibres based on the polymer blend with 25 wt% PEO
and varying amounts of BC, showed similar levels of PEO crystal-
linity (ca. 1.2 vol%, see Fig. 4a), however it was difficult to accurately
assess the PEO crystallinity in this case, due to the presence of over-
lapping peaks from the crystalline cellulose.
Fig. 4b shows polarized IR spectra of the aligned electrospun
PMMA with 25 wt% PEO. The absorption peak at 749 cm
21
of the
skeletal vibrational motion of the CH
2
group in PMMA
38
showed IR-
dichroism; the angle between the transition moment vector and the
chain axis (denoted a) has been reported to be 17u
33
. The IR-dichro-
ism; R 5A
jj
/A
H
51.228 (A
jj
,A
H
refer to the absorbance values for
light polarized parallel to and perpendicular the fibre axis)
39
. The
Herman’s orientation function (f) was calculated according to
40–42
:
f~1
23 cos2f
{1
~R{1
Rz2
:2 cot2aðÞz2
2 cot2aðÞ{1ð1Þ
Figure 3
|
(a) Stress-strain curves of aligned PMMA/PEO fibre mats containing different amounts of bacterial cellulose (BC) (3–5) compared to those of
pristine PMMA and PMMA fibres with 25 wt% PEO; SEM fibre fracture cross sections of (a) pristine PMMA, (b) PMMA with 25 wt% PEO, (d) PMMA
with 25 wt% PEO and 10 wt% BC; (e) TEM of 10 wt% BC in PMMA with 25 wt% PEO.
Table 2
|
Mechanical properties of electrospun composite fibres containing reinforcing cellulose crystals. Highest values are bold
Fibres formulation Properties
PMMA
100 [%]
PMMA/WC
96/4 [%]
22
PMMA/PEO
75/25 [%]
PMMA/PEO 1BC
75/25 14 [%]
PMMA/PEO 1BC
75/25 110 [%]
Toughness [MJ/m
3
] 0.13 60.04 0.16 60.04 17.1 61.4 1.09 60.16 0.61 60.03
Ultimate strength [MPa] 16.8 60.4 21.9 63.7 18.7 60.4 17.9 60.57 27.6 62.4
Young’s modulus [GPa] 1.49 60.29 1.93 60.29 0.72 60.05 0.91 60.02 1.95 60.11
Fibre diameter [mm] 1.74 60.35 2.70 60.58 1.96 60.32 1.94 60.36 1.97 60.41
Alignment std. dev. 14.3u14.7u12.5u13.1u13.8u
Rel. toughness [%] ref
1
120
1
113 053 1738 1369
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 4
where fis the angle between chain axis and the director (fibre axis).
The chain orientation fwas 0.081. PMMA/PEO (ca. 25 wt% PEO)
blend films required a draw ratio of 3 in order to achieve a similar
orientation factor
39
. Since this measurement was made on a fibre mat
with a non-perfect alignment of fibres (Fig. 1d and Table 1), the
orientation of the individual fibres must have been slightly higher.
However, the extensive increase in toughness of the electrospun
fibres compared to previous data for PMMA/PEO, the cast films
or the extruded bars, could be explained neither by the moderate
initial molecular alignment nor by the crystallinity of the fibres after
spinning.
Mechanisms for toughness improvement.The decrease in cross
sectional area by a factor of 7.6 (3.8 to 0.5 um
2
) can be explained
by considering that only a portion of the fibres undergo necking, i.e.
via the necking deformation zone shown in Fig. 2 (whereas an affine
deformation gives a reduction factor of only 3). The size of this
deformation zone at any given time during deformation was ca.
100–200 nm in the case of the electrospun fibres. This length can
be compared to the size of the PEO (
Mwof 600 kDa) and the PMMA
(
Mwof 410 kDa) molecules, which would be 5.0 and 1.0 mm long if
they were completely straight (0.154 nm bond length)
42,43
.Itis
therefore clear that the molecules could easily bridge the necking
zone in a continuous stretching event (alignment) as the
deformation zone propagated along the fibres during deformation.
However, a more realistic explanation would be to relate the end-to-
end distance of a random coil of the PEO polymer which in its coiled
state can be approximately 66 nm (r56
1/2
r
g
2
)
43
, based on a radius of
gyration (r
g
) of ca. 27 nm to the length scale of the deformation
44
.
This end-to-end distance is smaller than the deformation zone but,
since the electrospinning induces orientation of the polymer chains,
the true end-to-end distance in the direction of the fibre can be
expected to be significantly larger and to correspond to the initial
molecular orientation, which was equivalent to a tensile draw ratio of
3 (Fig. 4b)
39
. The end-to-end distance along the direction of the fibre
is therefore expected to cover about the same or a slightly larger
distance than the deformation zone during the necking. This
would allow a continuous stretching and energy build-up in the
deformation zone, which greatly contributed to the toughness of
the fibres. The molecules thus adopted an unprecedented degree of
orientation after the deformation zone, and this may resemble strain
hardening
45
and allow the necks to bear a higher load per cross–
sectional area than the undeformed parts of the fibres.
A possible contribution to the necking behaviour is the presence of
a thin ductile surface layer with a lower glass transition temperature
(T
g
), which prevented crack initiation by more facile reorientation of
molecules in the neck deformation zone. This surface layer of the
same molecular composition as the interior of the fibres emanates
from an enhanced molecular mobility in the proximity of the fibre
surface
46,47
. Evidence of a gradually decreasing T
g
with closer distance
to the surface of polystyrene was recently reported, wherein a sup-
pression of ca. 50uC (or more) was demonstrated 10–50 nm from the
polymer surface, depending on the measurement technique used
46,47
.
An equivalent suppression of the T
g
for the 75/25 PMMA/PEO blend
would result in a glass transition below room temperature, i.e. a
‘rubbery’ material at the surface of the fibres. To explore this hypo-
thesis, the toughest fibre formulation (75/25 PMMA/PEO) was elec-
trospun at the slowest feed–rate possible to generate uniform fibres
(10 mL/min) with approx. a doubled specific surface area. The
obtained fibres (diameter: 1.2 60.1 mm) showed a Young’s modulus
that was reduced by ca. 50% and with similar strength as the 2 mm
fibres. The lower modulus and more ductile fibres obtained therefore
served as an indication that a larger portion of rubbery surface layer
existed, assuming the existence of a rubbery surface layer of the same
thickness. However, the obtained fibres also showed a ca. 65%
decreased strain to failure, which was related to an increased molecu-
lar alignment for thinner fibres (as suggested by Papkov et al.)
6
,
which prohibited them from aligning further during the tensile test.
It is also clear that the probability of encountering defects at any
given time during deformation becomes smaller as the volume assoc-
iated with the deformation decreases, and that the mechanical prop-
erties of the entire mat are related to neck formation distributed in
thousands of fibres deforming simultaneously. An approximate cal-
culation shows that the deformation zone constitutes less than 1 ppm
of the entire fibre length. Hence, even when a neck propagates upon a
defect, the surrounding thousands of fibres can continue to stretch
under load, and absorb energy in the absence of defects. Based on the
above reasoning and the assumption that the non-necking portion of
the fibre did not deform at all, the maximum volume of the neck was
calculated to be 23% of the total fibre volume (based on the reduction
in fibre diameter via necking and the 150% strain at break). This
‘‘necked’’ volume of the fibres corresponded to 71% of the final fibre
Figure 4
|
WAXD diffractograms of the electrospun PMMA/PEO fibres with and without bacterial cellulose (a), IR spectrum for the toughest fibres
containing 25 wt% PEO (b) with inset IR-dichroism (polarized IR) for the CH
2
skeletal vibration peak in PMMA at 749 cm
21
.
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 5
length. Hence, if the prepared fibre formulation had allowed com-
plete propagation of the necks over the entire fibre length, the tough-
ness of the fibres would have increased even more, i.e. if the
remaining 77 vol% of the fibres had been allowed to form necks in
the absence of defects. An upper limit for the toughness of electro-
spun fibres of a PMMA/PEO blend with 25% PEO could then be
calculated to be approximately 75 MJ/m
3
, which is more than 300%
the observed toughness, and comparable to that of carbon steel
48
.
Conclusions
This work presents for the first time electrospinning as a method of
preparing ultra-toughened biocompatible fibre networks based on
blends of the inexpensive engineering polymers poly(methyl metha-
crylate) (PMMA) and polyethylene oxide (PEO). The key to the
toughening was to prepare a sufficiently ductile polymer blend,
which could undergo an extensive and energy-absorbing fibre neck-
ing phenomenon prior to fibre failure due to a suppressed glass
transition temperature at the surface of the fibres. A simultaneous
increase in toughness (27 times), modulus (10.7%) and tensile
strength (77%) was observed for fibres containing 12.5 wt% PEO,
25 wt% PEO, the brittle PMMA showed an increase in toughness of
more than two orders of magnitude. The toughness improvement
was accompanied by an increase in true fibre strength of ca. one
order of magnitude from 16.8 MPa to ca. 150 MPa, as verified
in in-situ electron micrographs during the tensile testing. Cast films
and extruded rods based on the same 25 wt% PEO blend showed
respectively 61% and 82% lower toughness than the electrospun
counterpart. It is suggested that molecular orientation at the
energy-absorbing neck requires an initial degree of molecular align-
ment (obtained via electrospinning) that is sufficiently high so that
the size of the polymer molecules in the direction of the fibre is of the
same order of magnitude as the deformation zone. Fibres conse-
quently break when the deformation zone reach a point where the
initial alignment is insufficient. This reasoning is consistent with our
findings that cast films and extruded rods, with a larger size and with
presumably little alignment, have toughness inferior to that of the
electrospun counterpart. To further investigate the necking ability of
the electrospun fibres, experiments were performed with the addi-
tion of highly crystalline cellulose fibres (as rod-like reinforcement)
to the PMMA/PEO solution prior to electrospinning. This resulted in
hybrid composite fibres that were stiffer and stronger, but also
showed dramatically reduced toughness due to an inability of the
rod-like cellulose to adapt to the extensive deformation inherent in
neck formation.
The reported phenomena are likely to change the views on how to
develop simultaneously strong and strain compliant polymer formu-
lations for thin-fibre applications, since the reported phenomena
become apparent only when dimensions approach scales close to
one hundredth of the width of a hair strand.
Methods
Preparation of fibre solutions.PMMA with a
Mwof 410 kDa (Alfa Aesar) and PEO
(Acros Organics) with a
Mwof 600 kDa were dissolved in DMF (BDH Prolabo
99.8%), or in DMF/BC containing suspensions (the details of the preparation of BC
with 72% crystalline content
30
are provided in the Supplementary Information S1).
The polymer solutions/suspensions were heated to 70uC and kept at this temperature
under constant stirring until homogeneous and completely transparent suspensions
were obtained (2 h). The solutions were electrospun within 24 h.
Electrospinning and collection of fibre mats.The fibre solutions were continuously
fed from a 5 mL solvent-resistant polypropylene syringe at a rate of 40 mL/min
(60.1%), via a PTFE tube, to a flat tip 18-gauge needle with internal diameter of 0.84
mm. The needle tip was positioned 240 mm vertically above the rotating collector,
and the electric field from the needle to the collecting surface was maintained at ca.
40–50 kV/m, where the more viscous solutions required a higher electric field. The
electrospun fibres were deposited on the rotating cylindrical collector/drum at 2000
rpm (collector diameter 550 mm). Thermo-gravimetric measurements were carried
out on all the samples to ensure that the fibre mats contained no residual low volatile
DMF solvent prior to tensile testing. Infrared (IR) spectroscopy was also performed
for all samples prior to further characterization, in order to ensure that there was no
solvent remaining inside the fibres (Supplementary Information S3).
Tensile testing of fibres.The tensile measurements were performed on a Deben
Microtest stage, modified for direct A/D converter readout, which increased the
resolution of the load cell (76 mN) and extensometer (167 nm). The measurements
were carried out at a strain rate of 0.5 mm/min (10% of sample length per minute) and
the tensile stress values were calculated by dividing the measured force by the cross
sectional area of the fibre mat. The thickness of the fibre mat was derived from the
area density (mass per unit area) of the electrospun fibre mat, as determined from the
weight and area of an adjacent fibre mat on the same collector. The fibre tensile testing
was performed according to the template transfer method (TTM)
22
. A layer of
fluorinated ethylene propylene release film (thickness 576 mm) was attached to the
collector, followed by a pre-cut aluminium foil template (thickness 530 mm) that
was firmly attached to the surface of the collector directly over the release film with
conductive copper tape (thickness 566 mm). The template had a cut-out window
(width: 10 mm, length: 5 mm) that acted as a gap where fibres could freely span. Fig. 5
shows an illustration of the TTM method.
The fibres were fixed on both sides of the window using alkoxy-ethyl-cyanoacrylate
(Loctite 460, Henkel AG & Co. KGaA, Germany) before the aluminium template was
moved to the tensile tester. The side panels of the pre-cut window were cut (without
touching the fibres) immediately before testing when the template had been firmly
mounted in the tensile stage
22
.
Preparation of cast films.The solution used for electrospinning was also used to cast
films inside petri dishes by evaporating the solvent at a temperature of 50uCin
vacuum for 2 days to ensure a complete evaporation of all the solvent. The films were
cut and mounted for tensile testing in the same manner as the electrospun fibres.
Figure 5
|
(a) Template transfer configuration, illustrating the aligned fibres on the rotating collector and the aluminium foil cutout template. (b)
Illustration of the fixed fibres on the template as mounted in the tensile stage during tensile testing.
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 6
Melt extrusion of bars.Powders of the same PMMA and PEO material were mixed in
a DSM Xplore 5 ml micro twin-screw extruder at a temperature of 185uC. A short
sleeve die was fitted with a 2 mm diameter that induced minimal molecular
orientation to the final bars, and this resulted in a final diameter of ca. 2.7 mm (due to
die swelling). These much larger specimens than the cast films and electrospun fibres
were tested in an Instron 5566 tensile testing machine with a specimen length of 40
mm, using the same strain rate as for all the other samples.
Microscopy.A Hitachi S-4800 cold–field–emission scanning electron microscope
(SEM) was used. A ca. 2 nm coating of platinum–palladium was sputtered onto the
surface of the samples (10 s at 80 mA) in a Cressington 208HR high-resolution
sputter. The in-situ SEM tensile tests were performed in a Hitachi T-100 charge
reduction SEM that allowed testing without sputtering in order not to alter the
properties of the material.
X–ray diffraction (XRD).XRD measurements were conducted on a PANalytical
X’Pert Pro diffractometer using Cu–Karadiation (45 kV, 35 mA) and a 1.00 arcmin
step size. Crystallinity was calculated from the ratio of the integrated intensity of the
crystalline peaks and the total intensity of the diffractograms (including the
amorphous component).
Thermal analysis.Thermal characteristics of the electrospun mats were assessed by
differential scanning calorimetry (DSC) using a Mettler Toledo DSC 1 with 40 mL
aluminium cups, and by thermal-gravimetric analysis (TGA) using a Mettler Toledo
TGA/DSC 1 with 70 mLAl
2
O
3
crucibles. The heating rate was 10uC/min and a
nitrogen gas flow rate of 10 ml/min was used.
Infra red (IR) spectroscopy.IR spectroscopy was performed on a Perkin-Elmer
Spectrum 2000 using a 1 cm
21
scan step equipped with a linear polarizer and a single
reflection attenuated total reflectance stage (ATR) MKII Golden Gate unit (Specac
Ltd., London, UK).
1. Luo, C. J., Stoyanov, S. D., Stride, E., Pelan, E. & Edirisinghe, M. Electrospinning
versus fibre production methods: from specifics to technological convergence.
Chem. Soc. Rev. 41, 4708–4735, DOI:10.1039/c2cs35083a (2012).
2. Hou, H. Q. et al. Electrospun polyacrylonitrile nanofibers containing a high
concentration of well-aligned multiwall carbon nanotubes. Chem. Mater. 17,
967–973, DOI:10.1021/Cm0484955 (2005).
3. Jiang, S., Hou, H., Greiner, A. & Agarwal, S. Tough and transparent nylon-6
electrospun nanofiber reinforced melamine-formaldehyde composites. ACS Appl
Mater Interfaces 4, 2597–2603, DOI:10.1021/am300286m (2012).
4. Bergshoef, M. M. & Vancso, G. J. Transparent nanocomposites with ultrathin,
electrospun nylon-4,6 fiber reinforcement. Adv. Mater. 11, 1362–1365,
DOI:10.1002/(Sici)1521-4095(199911)11:161362::Aid-Adma1362.3.0.Co;2-X
(1999).
5. Swart, M., Olsson, R. T., Hedenqvist, M. S. & Mallon, P. E. Organic-Inorganic
Hybrid Copolymer Fibers and Their Use in Silicone Laminate Composites.
Polym. Eng. Sci. 50, 2143–2152, DOI:10.1002/pen.21749 (2010).
6. Papkov, D. et al. Simultaneously Strong and Tough Ultrafine Continuous
Nanofibers. ACS Nano 7, 3324–3331, DOI:10.1021/nn400028p (2013).
7. Li, D., Wang, Y. L. & Xia, Y. N. Electrospinning of polymeric and ceramic
nanofibers as uniaxially aligned arrays. Nano Lett. 3, 1167–1171, DOI:10.1021/
Nl0344256 (2003).
8. Zhou, C., Chu, R., Wu, R. & Wu, Q. Electrospun Polyethylene Oxide/Cellulose
Nanocrystal Composite Nanofibrous Mats with Homogeneous and
Heterogeneous Microstructures. Biomacromolecules 12, 2617–2625,
DOI:10.1021/bm200401p (2011).
9. Demczyk, B. G. et al. Direct mechanical measurement of the tensile strength and
elastic modulus of multiwalled carbon nanotubes. Mat Sci Eng a-Struct 334,
173–178, DOI:10.1016/S0921-5093(01)01807-X (2002).
10. Sung, J. H., Kim, H. S., Jin, H. J., Choi, H. J. & Chin, I. J. Nanofibrous membranes
prepared by multiwalled carbon nanotube/poly(methyl methacrylate)
composites. Macromolecules 37, 9899–9902, DOI:10.1021/Ma048355g (2004).
11. Huang, C. B., Chen, S. L., Reneker, D. H., Lai, C. L. & Hou, H. Q. High-strength
mats from electrospun poly(p-phenylene biphenyltetracarboximide) nanofibers.
Adv. Mater. 18, 668, DOI:10.1002/adma.200501806 (2006).
12. Olsson, R. T. et al. Extraction of Microfibrils from Bacterial Cellulose Networks for
Electrospinning of Anisotropic Biohybrid Fiber Yarns. Macromolecules 43,
4201–4209, DOI:10.1021/ma100217q (2010).
13. Asran, A. S., Seydewitz, V. & Michler, G. H. Micromechanical properties and
ductile behavior of electrospun polystyrene nanofibers. J. Appl. Polym. Sci. 125,
1663–1673, DOI:10.1002/App.34847 (2012).
14. Tan, E. P. S. & Lim, C. T. Physical properties of a single polym eric nanofiber. Appl.
Phys. Lett. 84, 1603–1605, DOI:10.1063/1.1651643 (2004).
15. Lim, C. T., Tan, E. P. S. & Ng, S. Y. Effects of crystalline morphology on the tensile
properties of electrospun polymer nanofibers. Appl. Phys. Lett. 92, DOI:10.1063/
1.2857478 (2008).
16. Wong, S. C., Baji, A. & Leng, S. W. Effect of fiber diameter on tensile properties of
electrospun poly(epsilon-caprolactone). Polymer 49, 4713–4722, DOI:0.1016/
j.polymer.2008.08.022 (2008).
17. Arinstein, A., Burman, M., Gendelman, O. & Zussman, E. Effect of
supramolecular structure on polymer nanofibre elasticity. Nat Nanotechnol 2,
59–62, DOI:10.1038/nnano.2006.172 (2007).
18. Pai, C. L., Boyce, M. C. & Rutledge, G. C. Mechanical properties of individual
electrospun PA 6(3)T fibers and their variation with fiber diameter. Polymer 52,
2295–2301, DOI:10.1016/j.polymer.2011.03.041 (2011).
19. Chew, S. Y., Hufnagel, T. C., Lim, C. T. & Leong, K. W. Mechanical properties of
single electrospun drug-encapsulated nanofibres. Nanotechnology 17, 3880–3891,
DOI:10.1088/0957-4484/17/15/045 (2006).
20. Shin, M. K. et al. Size-dependent elastic modulus of single electroactive polymer
nanofibers. Appl. Phys. Lett. 89, DOI:10.1063/1.2402941 (2006).
21. Naraghi, M., Arshad, S. N. & Chasiotis, I. Molecular orientation and mechanical
property size effects in electrospun polyacrylonitrile nanofibers. Polymer 52,
1612–1618, DOI:10.1016/j.polymer.2011.02.013 (2011).
22. Andersson, R. L. et al. Micromechanical Tensile Testing of Cellulose-Reinforced
Electrospun Fibers Using a Template Transfer Method (TTM). J. Polym. Environ.
20, 967–975, DOI:10.1007/s10924-012-0486-6 (2012).
23. Allan, B. Closer to nature: new biomaterials and tissue engineering in
ophthalmology. Brit J Ophthalmol 83, 1235–1240, DOI:10.1136/bjo.83.11.1235
(1999).
24. Amon, M. & Menapace, R. Cellular invasion on hydrogel and poly(methyl
methacrylate) implants: An in vivo study. J. Cataract Refract. Surg. 17, 774–779,
DOI:10.1016/S0886-3350(13)80410-5 (1991).
25. Dugan, J. M., Gough, J. E. & Eichhorn, S. J. Bacterial cellulose scaffolds and
cellulose nanowhiskers for tissue engineering. Nanomedicine 8, 287–298 (2013).
26. Shi, W. & Han, C. C. Dynamic Competition between Crystallization and Phase
Separation at the Growth Interface of a PMMA/PEO Blend. Macromolecules 45,
336–346, DOI:10.1021/ma201940m (2012).
27. Li Xuan, W. Y. & Yang Yang. Studies of the crystallization behavior in the
crystalline/amorphous polymer blends: poly(ethylene oxide)/poly(methyl
methacrylate) and poly(ethylene oxide)/poly(vinyl acetate). Polym. Commun.
280–288 (1985).
28. Lodge, T. P., Wood, E. R. & Haley, J. C. Two calorimetric glass transitions do not
necessarily indicate immiscibility: The case of PEO/PMMA. J. Polym. Sci., Part B:
Polym. Phys. 44, 756–763, DOI:10.1002/polb.20735 (2006).
29. Schwahn, D., Pipich, V. & Richter, D. Composition and Long-Range Density
Fluctuations in PEO/PMMA Polymer Blends: A Result of Asymmetric
Component Mobility. Macromolecules 45, 2035–2049, DOI:10.1021/ma2019123
(2012).
30. Sacui, I. et al. Comparison of the properties of cellulose nanocrystals and cellulose
nanofibrils isolated from bacteria, tunicate, and wood processed using acid,
enzymatic, mechanical, and oxidative methods. ACS Appl Mater Interfaces 6,
6127–6138, DOI:10.1021/am500359f (2014).
31. Nishiyama, Y. Structure and properties of the cellulose microfibril. J. Wood Sci. 55,
241–249, DOI:10.1007/s10086-009-1029-1 (2009).
32. Iwamoto, S., Kai, W., Isogai, A. & Iwata, T. Elastic Modulus of Single Cellulose
Microfibrils from Tunicate Measured by Atomic Force Microscopy.
Biomacromolecules 10, 2571–2576, DOI:10.1021/bm900520n (2009).
33. Hepworth, D. G. & Bruce, D. M. A method of calculating the mechanical
properties of nanoscopic plant cell wall components from tissue properties.
J Mater Sci 35, 5861–5865, DOI:10.1023/a,1026716710498 (2000).
34. Mark, R. E. Cell wall mechanics of tracheids. (Yale University Press, 1967).
35. Kroonbatenburg, L. M. J., Kroon, J. & Northolt, M. G. Chain Modulus and
Intermolecular Hydrogen Bonding in Native and Regenerated Cellulose Fibres.
Polym. Commun. 27, 290–292 (1986).
36. Martuscelli, E., Silvestre, C., Addonizio, M. L. & Amelino, L. Phase-Structure and
Compatibility Studies in Poly(Ethylene Oxide) Poly(Methyl Methacrylate)
Blends. Makromol Chem 187, 1557–1571 (1986).
37. Lezcano-Gonzalez, I. et al. Chemical deactivation of Cu-SSZ-13 ammonia
selective catalytic reduction (NH3-SCR) systems. Applied Catalysis B:
Environmental 154–155, 339–349, DOI:10.1016/j.apcatb.2014.02.037 (2014).
38. Nagai, H. Infrared spectra of stereoregular polymethyl methacrylate. J. Appl.
Polym. Sci. 7, 1697–1714, DOI:10.1002/app.1963.070070512 (1963).
39. Zhao, Y., Jasse, B. & Monnerie, L. Orientation and relaxation in uniaxially
stretched poly(methyl methacrylate) poly(ethylene oxide) blends. Polymer 30,
1643–1650, DOI:10.1016/0032-3861(89)90324-8 (1989).
40. Hermans, P. H. & Platzek, P. Beitra¨ge zur Kenntnis des
Deformationsmechanismus und der Feinstruktur der Hydratzellulose. Kolloid-Z
88, 68–72 (1939).
41. Hermans, P. H. Contribution to the physics of cellulose fibres. (Elsevier Publishing
Company Inc., 1946).
42. Gedde, U. W. Polymer Physics. (Springer, 1995).
43. Debye, P. The intrinsic viscosity of polymer solutions. J. Chem. Phys 14, 636–639
(1946).
44. Branca, C. et al. Study of Conformational Properties of Poly(ethylene oxide) by
SANS and PCS Techniques. Phys. Scr. 67, 551 (2003).
45. Birley, A. W., Haworth, B. & Batchelor, J. Physics of Plastics: Processing, Properties,
and Materials Engineering. (Hanser Gardner Publications, 1992).
46. Ellison, C. J. & Torkelson, J. M. The distribution of glass-transition temperatures
in nanoscopically confined glass formers. Nat Mater 2, 695–700 (2003).
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 7
47. Ba¨umchen, O., McGraw, J. D., Forrest, J. A. & Dalnoki-Veress, K. Reduced Glass
Transition Temperatures in Thin Polymer Films: Surface Effect or Artifact? Phys.
Rev. Lett. 109, 055701 (2012).
48. Callister, W. D. & Rethwisch, D. G. Fundamentals of Materials Science and
Engineering: An Integrated Approach. (Wiley, 2012).
Acknowledgments
The Swedish International Development Cooperation Agency (SIDA) is acknowledged for
financial support. The authors also thank M. Salajkova for her assistance in the high
resolution SEM imaging of PMMA/PEO cross sections.
Author contributions
The project was planned and overseen by M.H., U.G., P.M. and R.O. The experimental
measurements were performed by R.A. as well as preparation of figures. Data analysis was
performed by V.S. All authors contributed to the manuscript preparation.
Additional information
Supplementary information accompanies this paper at http://www.nature.com/
scientificreports
Competing financial interests: The authors declare no competing financial interests.
How to cite this article: Andersson, R.L. et al. Micromechanics of ultra-toughened
electrospun PMMA/PEO fibres as revealed by
in-situ
tensile testing in an electron
microscope. Sci. Rep. 4, 6335; DOI:10.1038/srep06335 (2014).
This work is licensed under a Creative Commons Attribution-NonCommercial-
NoDerivs 4.0 International License. The images or other third party material in
this articleare included in the article’sCreative Commons license, unless indicated
otherwise in the credit line; if the material is not included under the Creative
Commons license, users will need to obtain permission from the license holder
in order to reproduce the material. To view a copy of this license, visit http://
creativecommons.org/licenses/by-nc-nd/4.0/
www.nature.com/scientificreports
SCIENTIFIC REPORTS | 4 : 6335 | DOI: 10.1038/srep06335 8
