C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

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Carbonaceous fillers for shape memory actuationof polyurethane composites by resistive heating

I. Sedat Gunes, Guillermo A. Jimenez1, Sadhan C. Jana*

Department of Polymer Engineering, The University of Akron, 250 South Forge Street, Akron, OH 44325-0301, United States

A R T I C L E I N F O

Article history:

Received 4 August 2008

Accepted 28 November 2008

Available online 11 December 2008

0008-6223/$ - see front matter � 2008 Elsevidoi:10.1016/j.carbon.2008.11.053

* Corresponding author: Fax: +1 330 258 2339E-mail address: janas@uakron.edu (S.C. Ja

1 Present address: Laboratory of Polymers

A B S T R A C T

The effectiveness of carbonaceous, electrically conductive fillers in shape memory actua-

tion of polyurethane composites by resistive heating was evaluated. Specifically, the depen-

dence of electrical resistivity on specimen temperature and imposed tensile strains

encountered in shape memory test cycles was determined for shape memory polyurethane

(SMPU) composites of carbon nanofiber (CNF), oxidized carbon nanofiber (ox-CNF), and car-

bon black (CB). The SMPU composites with crystalline soft segments were synthesized from

diphenylmethane di-isocyanate, 1,4-butanediol, and poly(caprolactone)diol in a low-shear

chaotic mixer and in an internal mixer. The materials synthesized in the chaotic mixer

showed higher soft segment crystallinity and lower electrical percolation threshold. A

reduction in soft segment crystallinity was observed in the presence of CNF and ox-CNF;

the reduction was smaller in the case of ox-CNF. Only the composites of CB showed pro-

nounced positive temperature coefficient (PTC) effects. The observed PTC effects bore a

close relationship with non-linear thermal expansion during heating. The composites of

CNF and ox-CNF did not show PTC effects due to low levels of soft segment crystallinity.

The resistivity of composites of CB increased by several orders of magnitude with imposed

tensile strain while composites of CNF and ox-CNF showed weak dependence on strain.

� 2008 Elsevier Ltd. All rights reserved.

1. Introduction

Freeman Dyson, a famous physicist and popular science wri-

ter envisioned in 1986 the concept of ‘space butterfly’, a new

generation of micro-spacecraft with morphing antennae or

wing-like structures [1]. The antennae can be stored in com-

pact form, readily deployed in flight, and their orientation

and length span can be easily controlled by electric current

generating devices. This vision is now close to reality due to

recent advances in the area of electrically triggered shape

memory polymers (SMP) and SMP nanocomposites [2]. SMPs

are a class of stimuli responsive materials, which recover the

original shapes from large deformation when subjected to

an external stimulus [3]. SMPs are commonly multi-phase

er Ltd. All rights reserved

.na).

(POLIUNA), School of Che

materials comprised of a fixed phase and a reversible phase.

The fixed phases are usually formed by thermally stable

‘cross-link’ points such as crystals, glassy domains, chain

entanglements, or chemical cross-links that prevent ‘free’

flow of surrounding polymer chains upon application of stress

even at elevated temperatures. The reversible phase, on the

other hand, undergoes deformation and shrinkage in a shape

memory cycle and is responsible for elasticity in the materials

[4]. A basic mechanism of shape recovery involves ‘shrinkage’

of oriented, extended chains of the reversible phase triggered,

for example, by melting or glass transition [5]. SMPs undergo-

ing shape recovery by shrinkage of oriented chains are classi-

fied as rubberlike SMP [6]. In rubberlike SMPs, the deformation

of chains and chain segments is preserved due to vitrification

.mistry, Universidad Nacional, Heredia, Costa Rica.

982 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

by glass transition or by crystallization. SMPs with semi-crys-

talline reversible phases possess certain advantages over the

glassy ones [7], e.g., shape recovery occurs in a narrow temper-

ature range due to melting of the crystalline part. This pro-

vides a relatively sharp and more complete shape recovery.

In addition, the semi-crystalline SMPs show higher toughness

due to the presence of crystalline soft segments. However, it is

imperative that SMPs with crystalline reversible phases should

possess adequate crystallinity in order to offer strong shape

memory performance.

The choice of particulate fillers in filled SMP systems may

aid in triggering SM function by mechanisms other than di-

rect application of heat. For example, SMP functions can be

actuated by application of light or electrical voltage if appro-

priate fillers are chosen, even though the matrix polymer

may not respond to such stimuli. This added functionality

can have profound importance in practical applications, such

as in aerospace industries. SMPs triggered by photon absorp-

tion are often limited to thin specimens. In view of this, SMPs

triggered by resistive heating can especially be useful in appli-

cations requiring remote actuation and the use of relatively

thicker specimens.

Shape memory actuation by resistive heating [8–10] has

been demonstrated for electrically conductive SMP compos-

ites containing both crystalline and glassy reversible phases.

In typical shape memory cycles, the composite specimens

are subjected to large deformations and large changes in tem-

perature. In view of this, the knowledge of temperature and

strain dependent electrical resistivity of SMP composites is

of central importance. A major challenge in the design and

functioning of SMP composite actuators based on semi-crys-

talline SMP is the positive temperature coefficient (PTC) of

resistivity, due to which otherwise electrically conductive

compounds transform into insulators as the temperature is

raised. This is often compounded by simultaneous melting

of the crystals, which in turn disrupts electrically conductive

networks. Consequently, further resistive heating with con-

tinued application of electrical voltage is not possible [11].

An avenue to alleviate PTC effect may be to use higher

amounts of conductive fillers, such that enough electrically

conductive networks survive after melting of the crystals.

However, it is known that higher filler content often reduce

the overall crystallinity and hence severely deteriorate the

shape memory properties of the composites [12].

Another critical aspect of shape memory actuation by

resistive heating is the dependence of electrical conductivity

on applied strain. Note that SMPs experience strains on the

order of several hundred percents both during deformation

and shape recovery [6,7]. It is imperative, therefore, that the

electrical conductivity of SMP composites should not change

much with large strains.

In this study, the issues of PTC effects and the relationship

between applied strain and electrical resistivity were investi-

gated by considering polyurethane SMP with semi-crystalline

reversible phase and three electrically conductive filler parti-

cles. A detailed evaluation of shape memory properties of

SMPU composites triggered by sensible heating was also car-

ried out and reported elsewhere [12,13]. In this paper, the rela-

tionships between soft segment crystallinity and the

electrical percolation behavior of SMPU composites filled with

carbon nanofiber (CNF), oxidized carbon nanofiber (ox-CNF),

and CB are first presented. The nature and origin of PTC ef-

fects, the effect of strain on electrical conductivity of the com-

posites, and their importance on shape recovery of SMP

composites triggered by resistive heating are then discussed.

2. Experimental

2.1. Materials

SMPU with 33% hard segment was synthesized from diph-

enylmethane diisocyanate (MDI, Bayer MaterialScience, Pitts-

burg, PA), 1,4-butanediol (BD, Avocado Organics, UK), and

polycaprolactonediol (PCL diol, Solvay Chemical, UK) of

molecular weight 4000 by mixing the ingredients in the stoi-

chiometric ratio of 6/5/1 by moles, respectively. The SMPU

thus produced contained a semi-crystalline soft segment

with melting point of approximately 45 �C. A tin catalyst,

DABCO T120 (Air Products, Allentown, PA) was used to expe-

dite chain extension reactions. Carbon nanofiber (CNF, Pyro-

graph III� PR-24-PS) and oxidized carbon nanofiber (ox-CNF,

Pyrograph III� PR-24-PS-ox), both vapor grown grade carbon

nanofibers, were obtained from Applied Sciences, Inc. (Cedar-

ville, OH) with mean diameter 60–200 nm and mean length

30–100 lm. The oxidized carbon nanofiber was produced by

controlled oxidation of CNF under air atmosphere, at 400–

500 �C [14]. High structure, conductive carbon black (CB,

Ketjenblack� EC 300 J) with pore volume of 0.310–0.345 ·10�3 m3/100 g [15], determined by dibutyl phthalate absorp-

tion was obtained from Akzo Nobel (Norcross, GA).

2.2. Preparation of composites

Composites were prepared separately in a chaotic mixer and a

commercial internal mixer. It was found earlier [16,17] that

composites prepared in the chaotic mixer shows electrical

percolation at much lower carbon nanofiber content than

materials prepared in the internal mixer when operated un-

der similar conditions of shear rate and temperature. This

was attributed to much less fiber damage in the chaotic mixer

and significant orientation of carbon nanofibers rendered by

the chaotic flow. Other researchers [18–20] reported lower per-

colation thresholds for composites of carbon black and ther-

moplastic polymers prepared in chaotic mixers. It was also

learned from previous work [17] that ox-CNF disperses well

in a polar polymer and consequently its percolation threshold

is higher than that of CNF, irrespective of the mixer used in

preparation of the composites. In view of these, it was ex-

pected that the composites of ox-CNF prepared in internal

mixer would require much higher filler concentration for

electrical percolation than those of CNF. To examine this is-

sue, four SMPU composites were selected in the present

study. Two sets of composites of CNF and ox-CNF were pre-

pared in the chaotic mixer. A second set of CNF composites

and one set of CB composites were prepared in the internal

mixer. Note that CB composites were not separately prepared

in the chaotic mixer.

The filler particles were dried overnight under vacuum at

120 �C to eliminate any absorbed moisture. Butanediol (BD)

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 983

and PCL diol were also dried overnight under vacuum at 45 �C;

nuclear magnetic resonance (NMR) and thermogravimetric

analysis (TGA) of dried BD and PCL diol specimens indicated

the absence of detectable residual water. All composites were

prepared using a two-step bulk polymerization method. The

prepolymer was synthesized from MDI and PCL diol and

chain extended with BD in the presence of tin catalyst. The

prepolymer was synthesized by allowing MDI and PCL diol

to react at 80 �C for 2 h under nitrogen with mechanical stir-

ring. The composites of CB and CNF with 1, 3, and 5 wt% fill-

ers were prepared in Brabender Plasticorder internal mixer

(model EPL 7752). For this purpose, the chain extension reac-

tion of prepolymer with BD was carried out at 110 �C for 2 min

in the presence of 9.8 · 10�4 mol/L tin catalyst. Filler particles

were then mixed with chain extended polymer for another

5 min with a set temperature of 140 �C. The temperature in

the mixer rose to 165 �C with the addition of filler particles

and then stabilized at 160 �C. The ox-CNF composites and a

second set of CNF composites were prepared in a chaotic mix-

er with mixing chamber volume of 70 cm3 and a gap of

12.5 mm between two circular rotors [21–23]. In this case,

approximately 80 g of molten prepolymer was hand-mixed

at 60 �C with proper amounts of CNF or ox-CNF followed by

mixing with 6.5 g of BD. The mixture was poured in the cha-

otic mixer preheated at 110 �C and mixed for 5 min. The circu-

lar rotors with a diameter of 12.7 mm were co-rotated in a

sinusoidal fashion (Eq. (1)) at peak angular speed (X, Eq. (1))

of 65 rpm, which generated a peak shear rate of 9.5 and

5.4 s�1 at the rotor surface and at the mixing chamber walls.

The time-averaged mean shear rate was 3.8 s�1. In Eq. (1),

X1 and X2 represent the angular speeds of rotors 1 and 2,

respectively, t is time, and T is the time-period

X1 ¼ X 1þ cos2ptT

� �; X2 ¼ X 1� cos

2ptT

� �; 0 6 t 6 T ð1Þ

The angular displacement of rotors in a time-period, h (�XT)

was 8p radians in order to produce wide-spread chaotic mix-

ing [22]. After synthesis, the composites were compression

molded at a pressure of 25 MPa and temperature of 220 �C into

specimens of 0.5 mm thickness. The total time for compres-

sion molding was 3 min, which prevented thermal degrada-

tion of the materials. The molded specimens were cold

compressed at room temperature for 15 min at a pressure of

25 MPa to produce flat, smooth specimen.

2.3. Characterization

The electrical resistivity of the composites was measured

using a four-probe Keithley 8009 resistivity tester and a Keith-

ley 487 picoammeter/voltage source obtained from Keithley

(Cleveland, OH). The dependence of electrical resistivity on

temperature was determined by measuring the electrical

resistivity while heating the composite specimens to desired

temperatures in a compression molder at a pressure of

5 MPa. It was confirmed that a minimum of 5 MPa pressure

was required for good mechanical contact. This was deter-

mined by measuring resistivity at various pressures up to

25 MPa. The susceptibility of electrical resistivity to imposed

tensile strain was determined by subjecting the composite

specimens to desired tensile strain at 60 �C in a tensile tester

(Instron 4204, Norwood, MA) fitted with a heating chamber.

The electrical resistivity of the specimen was simultaneously

measured. A layer of silver paste was applied on the compos-

ite specimens to obtain good electrical contact between the

specimen and the electrodes. Note that the conductive com-

posite specimens were stretched in a tensile tester at 60 �Cto produce specimens for shape memory tests.

The quality of filler dispersion in composite specimens

was examined by electron microscopy techniques. Approxi-

mately 70 nm thick specimens were microtomed under

cryogenic conditions using Reichert Ultracut S/FC S ultrami-

crotome (Leica, Germany) for transmission electron micros-

copy (TEM) using TEM device JEM-1200EXII (JEOL, Japan) at

120 kV. The scanning electron microscopy (SEM) images were

taken with SEM S-2150 microscope (Hitachi, Japan) at 20 kV

after sputter coating the cryogenically fractured specimens

with silver using a K575x sputter coater (Emitech, UK) under

argon atmosphere. The crystallinity and glass transition

temperature (Tg) of the soft segment phase were determined

using a differential scanning calorimetry (DSC) device, TA

instruments DSC-29210 (New Castle, DE) under nitrogen

atmosphere at a heating rate of 10 �C per minute in the

range �100 to 230 �C. During the first thermal scan, the spec-

imens were heated at a scanning rate of 10 �C/min to 230 �Cand then quenched to �100 �C at an average cooling rate of

70 �C/min using liquid nitrogen. The sample was subjected

to second thermal scan over a temperature range of �100

to 230 �C with a heating rate of 10 �C/min. The first thermal

scan was relatively more representative of the actual shape

memory testing cycle. The second thermal scan was per-

formed to gain further insight on crystallization behavior

of the composites. The DSC traces from the first and second

thermal scans were used to obtain the heat of fusion. The

percentage of crystallinity was determined by comparing

the heat of fusion obtained from DSC with the heat of fusion

of 100% crystalline PCL diol, 136 J/g [24]. PCL diol and BD

samples were examined for residual water after drying un-

der vacuum using TGA and NMR. A TGA equipment (TA

Instruments 2050, New Castle, DE) was used for this pur-

pose. A typical specimen of 6 mg weight was subjected to

analysis under nitrogen at a heating rate of 5�C/min, over

a temperature interval of 25–250 �C. The 1H NMR spectra

were recorded at room temperature with a Varian Mercury

300 MHz spectrometer (Varian Inc., Palo Alto, CA) with di-

methyl sulfoxide as the solvent. Thermal stability of poly-

urethane specimens was investigated using TGA with

typical specimens weighing approximately 6 mg subjected

to thermal scans from room temperature up to 1000 �C, un-

der nitrogen atmosphere, and at a scan rate of 20 �C/min.

The temperature at the onset of thermal degradation or 5%

mass loss, T1, was estimated from the specimen mass versus

temperature curves. The temperature at the maximum loss

rate, T2, was obtained from the plot of the first derivative

of the mass loss versus temperature curves. It was suggested

that the thermal degradation of PU occurred in two consec-

utive stages [25]. First, the degradation of hard segments

took place by dissociation of urethane groups to the initial

alcohol and isocyanate groups [25]. Further heating caused

depolymerization and degradation of soft segments [25].

984 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

Hence, the temperature T1 represented degradation behavior

of the hard segments, whereas the temperature T2 repre-

sented degradation of the soft segments. The molecular

weights (MW) of polymers were determined with gel perme-

ation chromatography method (GPC) using Waters 510 sys-

tem (Milford, MA) with triple detection scheme and a

polystyrene standard. The polydispersity index of the sam-

ple was determined from the ratio of weight average molec-

ular weight to the number average molecular weight. The

extent of hydrogen bonding in composite specimens, cast

as films from solution in dimethylformamide, was deter-

mined using Fourier-transform infrared spectroscopy

(FT-IR). For this purpose, the hydrogen bonding index, R

was determined as the ratio of areas of deconvoluted peaks

due to hydrogen bonded (�1700 cm�1) and free (�1730 cm�1)

carbonyl groups that appeared in infrared spectrum of

SMPU. The FT-IR spectra were recorded using Nicolet 4700

FT-IR Spectrometer with TGS Detector (Thermo Scientific,

Waltham, MA) in micro-attenuated total reflectance (micro-

ATR) mode. A micro-ATR fixture with a diamond crystal

and single bounce (SensIR Technologies, DuraScope, Dan-

bury, CT) was used for this purpose and scans were taken

at a resolution of 4 cm�1. The coefficient of thermal expan-

sion (CTE) of specimens was determined using a PerkinElmer

Pyris Diamond Dynamic Mechanical Analyzer (DMA, Wal-

tham, MA) in F-control mode. Note that the DMA fixtures

also underwent expansion and contraction, respectively,

during heating and cooling steps. The extent of such expan-

sion and contraction was determined by calibrating the DMA

set up using strips of copper with 99.999% purity. The heat-

ing rate was kept at 5 �C/min and at least 3 different repli-

cates for each condition were tested over the temperature

range of �60 to 120 �C. It was found that the CTE of DMA fix-

ture was 7.6 · 10�6/�C with a standard deviation of about 15%

[26]. The experimental data on composite specimens were

corrected using this value. It was reported elsewhere [26]

that the CTE of composite specimens included in this work

was approximately an order of magnitude higher than that

of the DMA fixture. The specimens were annealed at 90 �Cfor 48 h under vacuum before any testing was done. This

eliminated any potential chain orientation originated from

compression molding of the specimen. It was recognized

that residual chain orientation would have triggered shape

recovery as the test temperature approached the crystalline

melting region. Therefore, the values reported in this work

were not influenced by shape recovery of oriented chains.

Also note that the annealing temperature of 90 �C was se-

lected well below the onset temperature of thermal degrada-

tion of SMPU, �300 �C under nitrogen (Section 3.2). It was

also found that annealing under vacuum did not influence

the thermal and mechanical properties of the specimens.

The nature of surface functional groups on CB particles

was determined using X-ray photoelectron spectroscopy

(XPS) method. The XPS spectra were obtained using a Kratos

Model ES3000 spectrometer (Manchester, UK) under high

vacuum conditions with a pressure of 10�8 Torr, an alumi-

num anode, and a resolution of 1 eV. The assignment of

peak locations and corresponding fitting of XPS spectra were

performed with a curve fitting and data analysis software,

Fityk 0.7.7 (http://www.unipress.waw.pl/fityk).

3. Results and discussion

3.1. Morphology and surface characteristics of fillers

The morphology, graphitic content, and chemical nature of

surfaces of graphitic carbonaceous fillers bear close relation-

ship with the values of electrical resistivity. Typical resistivity

value of conductive CB is 0.5 X cm [27], carbon fiber (CF) is

0.001 X cm [28], multi-wall carbon nanotubes (MWCNT) is

50–100 · 10�6 X cm [29], and single-wall carbon nanotubes

[30] and graphite single crystals [31] is 40 · 10�6 X cm, in

increasing order of graphitic content. The resistivity of CNF

used in this work was reported to be 55 · 10�6 X cm by the

supplier. The single fiber resistivity data on ox-CNF was not

available, although it was expected to be lower than that of

CNF due to the presence of non-conductive oxygen contain-

ing functional groups on the fiber surfaces [17]. The single

fibers of CNF and ox-CNF differed in morphology as seen from

the TEM images in Fig. 1. Fig. 1a presents the cross-sectional

view of CNF which bears resemblance to the typical stacked

morphology of MWCNT [32]. It is evident from Fig. 1a that

CNF consisted of a number of concentric tubes, although like

MWCNT, many of them ceased to grow up to the very end of

the single fiber. A carbon nanotube is seen on the CNF surface

in Fig. 1a, purportedly grown from an occasional catalyst par-

ticle left on the nanofiber surface. Unlike the smooth surfaces

of CNF in Fig. 1a, the surfaces of ox-CNF particles contained

imperfections, as revealed from Fig. 1b. These imperfections

were generated at the time of oxidation of the nanofibers

and conversion of graphitic carbon into polar, organic func-

tional groups as was observed in the case of oxidized

MWCNTs by other researchers [33]. The conversion of gra-

phitic carbon in turn caused an increase of electrical resistiv-

ity of the oxidized fibers.

The polar, organic functional groups, such as ether, ester,

and carboxylic acid groups were identified on ox-CNF sur-

faces using XPS [17]. It was found that these polar functional

groups significantly improved the dispersion of ox-CNF parti-

cles in polar polymers, such as polyurethanes and polymeth-

ylmethacrylate (PMMA) [17]. In view of this, the CB particles

were also analyzed by XPS to determine the presence of polar,

organic functional groups on particle surfaces. The XPS spec-

tra of CB particles presented in Fig. 2 indicate the presence of

significant amounts of oxygen, apparently contained in polar

functional groups, such as ester (–COOR) and carboxylic acid

(–COOH); this is in accordance with the previous observations

involving various grades of CB particles [34]. The XPS scan

presented in Fig. 2 shows four peaks. Two peaks are due to

ionization of electrons located in 1s core levels of oxygen

and carbon and are identified as O1s and C1s, respectively

[35]. Two additional peaks originating from the secondary

(Auger) ionization of electrons in oxygen and carbon are ob-

served in the survey scan as C KVV and O KLL peaks [34].

The O1s and C1s peaks were analyzed further to characterize

the chemical nature of CB particle surfaces. The resolution of

C1s peak of CB (Fig. 3) revealed a peak at 288 eV assigned to

C@O containing groups [36] and another at about 284.2 eV as-

signed to graphitic carbon [17]. The ratio of carbon to oxygen

atoms on the CB particle surface was determined from the

areas under the curves for C1s and O1s peaks in Fig. 2 and

Original Graphitic

carbon

COOH COOR

Fig. 3 – The C1s region of CB XPS spectrum resolved into two

peaks.

Fig. 1 – (a) TEM image of CNF. The stacked morphology

typical for MWNT is evident. Note also the carbon nanotube

on the surface, purportedly grown from an occasional

catalyst particle left on the CNF surface. A fracture site

indicated also by an arrow is observed. (b) TEM image of ox-

CNF. The typical stacked morphology and imperfections

resulting from the surface oxidation (darker sites on the

fiber surface) are observed.

Fig. 2 – XPS spectrum of CB. O1s and C1s represent the

photoionization peaks originated from the electrons located

in 1s core levels of oxygen and carbon. Two more peaks

originated from the secondary (Auger) ionization of

electrons in oxygen and carbon is also identified as C KVV

and O KLL.

Table 1 – Atomic ratios of surface carbon atoms (areaunder the peak of C1s) to surface oxygen (area under thepeak of O1s) in the fillers determined with XPS. Data forCNF and ox-CNF were taken from a previous report [17].

Filler (C/O) ratio

CNF 66.7

Ox-CNF 7.7

CB 19.8

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 985

is presented in Table 1. It is apparent from Table 1 that of the

carbonaceous fillers used in this work, ox-CNF contained the

highest fractions of oxygen in the form of polar functional

groups, which in turn improved dispersion of ox-CNF in poly-

urethanes as reported previously [17].

3.2. Effects of fillers on soft segment crystallinity of thecomposites

The soft segment formed the reversible phase of SMPUs and

its crystallinity was of crucial importance for shape memory

performance. The soft segment crystallinity of SMPU and

glass transition temperature were determined by DSC and it

was found that the SMPU prepared in the chaotic mixer pro-

vided higher soft segment crystallinity, �30% and lower Tg

(�55 �C) compared to �20% crystallinity and �49 �C Tg for

materials prepared in Brabender internal mixer. The change

of soft segment crystallinity as function of filler content is

presented in Fig. 4. It was recently reported [21] that chaotic

mixers can be used effectively to obtain higher molecular

weight polyurethanes than commercial internal mixers when

operated under similar conditions of temperature and mean

shear rate. In the present study, the crystalline soft segment

originated from PCL diol. Therefore, higher soft segment crys-

tallinity should come from higher degree of phase separation

of the domains containing PCL diol. Note that hydrogen bond-

ing between the ester carbonyl groups of PCL diol and –NH

groups of urethane linkages promotes mixing of the soft

and hard segments and consequently may cause a reduction

Fig. 4 – The soft segment crystallinity of SMPU composites

determined with DSC (1st thermal scan). The composites

prepared in Brabender internal mixer and in chaotic mixer

are represented, respectively, by (B) and (C).

986 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

of the soft segment crystallinity. In view of this, the areas un-

der the peaks of bonded (AHCO) and free (ACO) carbonyl

absorption in FT-IR spectra of the composites were monitored

to determine the extent of hydrogen bonding between the

carbonyl groups in ester and the –NH groups of urethane link-

ages. It was found that the value of AHCO/ACO � 0.61 for the

neat SMPU prepared in chaotic mixer compared to �0.52 for

the materials prepared in the internal mixer. This indicates

that higher degree of phase separation occurred in materials

prepared in the chaotic mixer. The higher degree of phase

separation in materials prepared in the chaotic mixer ex-

plains the lower values of Tg of the soft segments observed

in Fig. 5 as highlighted by other researchers [37].

The representative FT-IR spectra of pristine SMPU and

CNF/SMPU composites with 5 wt% filler content, prepared

both in chaotic and internal mixers, are presented in Fig. 6a.

Note that carbonyl groups were present both in hard seg-

ments in the form of urethane linkages and in the soft seg-

Fig. 5 – Glass transition temperature of the soft segment in

CNF/SMPU composites prepared in internal (B) and chaotic

mixers (C).

ments in the form of ester linkages. Hence, it was difficult

to relate the value of hydrogen bonding index, R with the de-

gree of phase separation [38]. We observed from the data pre-

sented in Fig. 6b that the ratio of the area under the hydrogen

bonded C@O peak (1700 cm�1) to that of the free C@O peak

(1730 cm�1) in pristine SMPU and SMPU composites was not

significantly influenced by the presence of CNF. The same

was observed in cases of ox-CNF and CB. The DSC traces in

Fig. 7 indicate almost no changes in heat of fusion corre-

sponding to melting of the hard domains at around 225 �C.

Hence, it is conceivable that the presence of fillers exerted al-

most no impact on hard domains. The representative DSC

traces of CNF/SMPU composites presented in Fig. 7 show,

however, a remarkable reduction in heat of fusion corre-

sponding to melting of soft segments at around 40 �C.

In order to ensure that thermal degradation did not occur

during shape memory testing, the thermal degradation tem-

peratures (T1 and T2) of composite specimens CNF/SMPU

and ox-CNF/SMPU prepared in chaotic mixers were evaluated.

These results are presented in Fig. 8. It is seen that T1 and T2

followed similar trend, i.e., reduced in the presence of 1 wt%

filler and then increased with the addition of more filler. Sim-

ilar observations were reported for MWCNT/PU system [39]. It

is noted that MWCNT, CNF, and ox-CNF all enhance thermal

conductivity of SMPU [39], thus facilitate rapid heating of

the polymer, and hence expedite thermal degradation at low-

er filler contents. At higher loadings, however, MWCNT, CNF,

or ox-CNF produces a delay in thermal degradation, purport-

edly by producing barrier layers to degradation products [40].

The higher degree of phase separation of SMPU prepared

in chaotic mixer can be attributed to higher MW of the poly-

mer [41] – the number average MW of SMPU prepared in cha-

otic mixer was approximately 57,000 g/mol compared to

25,000 g/mol for SMPU prepared in the internal mixer. Similar

trend was observed in an earlier study on polyurethane syn-

thesis in chaotic mixers [21] and was attributed to better mix-

ing of the chain extender and the prepolymer in the chaotic

mixer. A few words are in order on the effectiveness of the

chaotic mixers in polyurethane polymerization. The rate of

condensation polymerization reactions, such as polyurethane

polymerizations, is known to be highly influenced by the rate

of interfacial area generation between the reactants [42]. In

this context, the mixing scheme in the polymerization reactor

can play a decisive role in determining molecular weight and

the final properties of the resulting polymers [43]. In our pre-

vious studies [21,44], we exploited the exponential interfacial

area growth feature of chaotic mixing and obtained higher

molecular weight polyurethanes with improved mechanical

and thermal properties. We also determined that the mixer

should operate under globally chaotic conditions to obtain

high molecular weight [21]. In view of these, we operated

the chaotic mixer at the same conditions as described in Refs.

[21,44].

Let us turn our attention to the effects of fillers on soft seg-

ment crystallinity. It was recently reported that the presence

of organoclay, nanometer size silicon carbide (SiC), CNF, and

CB all reduced the soft segment crystallinity in PCL diol-based

SMPU nanocomposites and composites prepared in an inter-

nal mixer [12,45]. However, the reduction was more signifi-

cant in the presence of CNF and SiC and it was observed

Wavenumber (cm-1)

1000 2000 3000 4000

Abs

orba

nce

0.0

0.2

0.4

0.6

0.8

1.0

1.2

pristineCNF1CNF3CNF5

Wavenumber (cm-1)

1000 2000 3000 4000

Abs

orba

nce

0.0

0.2

0.4

0.6

0.8

1.0

1.2

pristineCNF1CNF3CNF5

a

Fig. 6 – (a) Representative FT-IR spectra of CNF/SMPU composites prepared in internal (B) mixer. The number following ‘CNF’

indicates the filler content in the composite as weight percent. (b) Expanded view of free (1730 cm�1) and hydrogen bonded

(1700 cm�1) carbonyl stretching bands in SMPU and CNF/SMPU composites prepared in internal mixer.

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 987

that the filler shape, the number of filler particles per unit vol-

ume, and the nature of organic, polar functional groups on

the surface of filler particles, all play crucial roles in determin-

ing the soft segment crystallinity. In the present work, the vol-

umetric filler content in composites varied between 2 and 3%

due to differences in densities of the respective fillers, e.g.,

1.95 g/cm3 for CNF and ox-CNF [16] and 1.8 g/cm3 for CB [15]

even though the same weight percent of filler particles were

Fig. 7 – DSC heating scans of SMPU and CNF/SMPU

composites prepared in internal mixer.

a

b

Fig. 8 – Thermal stability of CNF/SMPU and ox-CNF/SMPU

composites prepared in chaotic mixer determined by TGA.

(a) The temperature at the onset of thermal degradation (T1).

(b) The temperature at the maximum rate of mass loss (T2).

Table 2 – Crystallinity of PCL diol composites with 10 wt%filler content.

Crystallinity (%)

1st scan 2nd scan

Neat PCL diol 58 53

CNF composite 57 54

Ox-CNF composite 59 54

CB composite 58 52

988 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

used. The soft segment crystallinity of SMPU reduced sub-

stantially in the presence of CNF and ox-CNF shown in

Fig. 4 irrespective of the mixer used. In comparison, CB parti-

cles with characteristic aggregate size of 0.4–1.0 lm [15]

caused only marginal changes in crystallinity. Note that the

typical aggregate size of CB particles was about two orders

of magnitude higher than the size of typical crystalline lamel-

lae �10 nm [46]. The more dramatic decrease in soft segment

crystallinity in the presence of CNF and ox-CNF can be attrib-

uted to their particle size and geometry, e.g., the characteris-

tic diameters of CNF and ox-CNF, 60–200 nm, were of the

same order as the size of the crystalline lamellae, which cou-

pled with their irregular, bundled morphology may have pre-

vented close packing and crystallization of the soft segment

[12].

The relatively higher soft segment crystallinity seen for

composites of ox-CNF (Fig. 4) can be explained as follows.

Note that the data in Fig. 4 were obtained under quiescent

conditions. A major difference between CNF and ox-CNF

was the chemical nature of their surfaces; a minor difference

might be the amorphous carbon content [17]. It was reported

earlier [17] that ox-CNF surface contained several polar organ-

ic groups (see also Table 1). Accordingly, the interfacial inter-

action between SMPU and ox-CNF was stronger due to the

presence of polar groups in both of them. It is still not clear

if additional polarity of ox-CNF surfaces had any bearing on

higher soft segment crystallinity as observed in Fig. 4.

A set of composites of CNF and ox-CNF were prepared by

mixing them with neat PCL diol. The objective was to deter-

mine if nanofiber surface polarity would also affect the crys-

tallinity of PCL diol molecules. The composites were prepared

by mixing the fillers with neat PCL diol in Brabender internal

mixer at 100 rpm for 5 min at 60 �C. The filler loading in these

composites was 10 wt%. The crystallinity data of the compos-

ites obtained during first and second scans are given in Table

2. It is seen that PCL diol crystallinity decreased only slightly

in the presence of the fillers and that the differences in polar-

ity of CNF and ox-CNF surfaces did not have any effect on PCL

diol crystallinity. In view of this, we invoked the differences in

the rate of crystallization as a possible reason why the same

nanofibers exerted different influence on crystallinity of com-

posites of SMPU and PCL diol. It is known that PCL diol is a

fast crystallizing polymer [47], whereas SMPU synthesized

from PCL diol undergoes slow crystallization [48].

A survey of literature (see Table 3) on polymer crystalliza-

tion in the presence of different fibers suggests that the rate

filler wt%

soft

seg

men

t cr

ysta

llini

ty (

%)

Fig. 9 – The soft segment crystallinity of CNF and ox-CNF

composites determined with DSC (2nd thermal scan). The

composites were prepared in chaotic mixer.

Table 3 – Comparison of the effects of substrate surface properties on crystallization of semi-crystalline polymers.

Polymer Substrate Observation onrate of crystallization

Significant effectof surface properties

of substrates oncrystallization

References

PE Aluminum oxide, PET Fast crystallizing No [49]

Nylon Fast crystallizing No [49]

PP Fast crystallizing No [49]

PP Aluminum oxide, copper oxide, PET, PTFE, FEP-PTFE Fast crystallizing No [50]

PP 43 pairs of substrates and crystallizing polymers Fast crystallizing No [51]

PEO Fast crystallizing No [51]

PCL Fast crystallizing No [51]

PB-1 Fast crystallizing No [51]

PP PET, PA 6,6 Fast crystallizing No [52]

HDPE Glass fibers with and without silane coupling agents Fast crystallizing No [53]

HDPE 3 types of carbon fibers (epoxy sized, unsized, and sized

and thermally treated)

Fast crystallizing No [54]

PCL diol Organoclay, CNF, SiC, CB Fast crystallizing No [12]

PCL diol CNF, ox-CNF, CB Fast crystallizing No This work

PPS 3 types of unsized carbon fibers with different

surface energies

Slow crystallizing Yes [55]

PPS Sized and unsized carbon fibers (type AS 4),

graphitized carbon fiber, sized and unsized

Kevlar aramid fibers, ‘PPS-compatible’ sized glass fibers

Slow crystallizing Yes [56]

Poly(ethernitrile) Pristine and oxidized carbon fibers Slow crystallizing Yes [57]

SMPU CNF, ox-CNF, CB Slow crystallizing Yes This work

SMPU Organoclay, CNF, SiC, CB Slow crystallizing Yes [12]

Abbreviations: PE: polyethylene, PP: polypropylene, PEO: poly(ethylene oxide), PCL: polycaprolactone, PB-1: poly(butane-1), HDPE: high density

polyethylene, PCL diol: polycaprolactone diol, PPS: poly(phenylene sulfide), SMPU: shape memory polyurethane, PET: poly(ethylene terepht-

alate), PTFE: poly(tetrafluoroethylene), FEP-PTFE: fluorinated ethylene-propylene copolymer, PA 66: polyamide 6,6, CNF: carbon nanofiber,

ox-CNF: oxidized carbon nanofiber, CB: carbon black, SiC: silicon carbide.

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 989

for fast crystallizing polymers is not affected by the presence

of fibers [49–54]. On the other hand, crystallization of slow

crystallizing polymers is significantly affected in the presence

of fibers [55–57]; the crystallinity significantly increases if the

polymer and the fiber surface are both polar or both non-po-

lar. Besides, polar–polar interactions between polar polymers

and polar adsorbents have been shown to significantly aug-

ment the adsorption of polar polymers onto the adsorbent

surface [58]. Polar–polar interactions at ox-CNF surfaces with

SMPU may have promoted adsorption of polymer chains on

the fiber surface, which in turn induced favorable chain con-

formations and hence higher rates of nucleation and crystal-

lization [59]. To verify this, crystallinity of SMPU composites

from a second DSC scan was determined. First the specimens

were heated from room temperature to 230 �C at a scanning

rate of 10 �C/min so as to melt the crystals and then quenched

to �100 �C at an average cooling rate of 70�C/min using liquid

nitrogen. In this manner, the soft segment of SMPU crystal-

lized in the presence of CNF and ox-CNF under identical ther-

mal gradient and the crystallinity determined in the second

thermal scan would reflect the specific influence of the fibers.

In the second thermal scan, the specimens were heated over

a temperature range of �100 to 230 �C with a heating rate of

10 �C/min. The crystallinity data thus obtained are presented

in Fig. 9, and show the same trend as in Fig. 4, i.e., crystallinity

in composites of ox-CNF were higher than in composites of

CNF. A slightly higher crystallinity in the composite of 1 wt%

ox-CNF compared to neat SMPU is also noted in Fig. 9.

3.3. Electrical conductivity of SMPU and PCL diolcomposites at room temperature

The volume electrical conductivity (r) as function of filler

weight fraction of SMPU composites is presented in Fig. 10.

It is seen that electrical percolation occurred at low filler con-

tent � 4 wt%, r � 10�5 S/cm, in composites of CNF prepared in

Fig. 10 – Electrical percolation in SMPU composites.

Composites prepared in Brabender internal mixer and in

chaotic mixer are represented, respectively, by (B) and (C).

1.00E-12

1.00E-09

1.00E-06

1.00E-03

1.00E+00

Fig. 11 – Electrical percolation in PCL diol composites.

990 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

the chaotic mixer. In comparison, the composites prepared in

Brabender internal mixer remained insulators, r � 10�11 S/cm

at up to 7 wt% CNF content. The low percolation threshold of

materials prepared in the chaotic mixer can be attributed to

much less attrition of high aspect ratio fibers and orientation

of the fibers rendered by the chaotic flow [17]. Recall that

higher aspect ratio fillers form conductive networks at lower

volume fraction and provide higher electrical conductivity

[60]. Also recall that electrical conduction in polymer compos-

ites is usually dominated by the inter-particle electron tun-

neling [61].

It is also evident from Fig. 10 that percolation threshold for

composites of ox-CNF prepared in the chaotic mixer is higher

�5 wt% with r � 10�5 S/cm. A similar observation was made

in the case of composites of poly(methylmethacrylate) [17]

and was attributed to better dispersion of ox-CNF in the poly-

mer. The composites of CB prepared in Brabender internal

mixer also showed a percolation threshold of �4 wt%, which

is close to data reported in literature [62] for CB/thermoplastic

polyurethane composites. The percolation threshold of CB is

usually low in multi-phase polymers, immiscible polymer

blends [20], and semi-crystalline polymers [63]. The lower per-

colation threshold in semi-crystalline polymers is attributed

to exclusion of conductive particles from the crystalline do-

mains, which in turn results in uneven distribution of con-

ductive particles and produces conductive networks at

lower filler loadings. The same trend was observed for electri-

cal percolation in composites of PCL diol as presented in

Fig. 11. Note that PCL diol is also a semi-crystalline polymer

with crystallinity of about 60% as determined from DSC and

better dispersion of ox-CNF in PCL diol matrix as in the case

of composites of SMPU resulted in higher percolation

threshold.

3.4. PTC effect

PTC effect in conductive polymer composites has been an

area of active research since the initial observation of the

phenomenon [64]. A critical review of experimental observa-

tions including the influence of different matrices and electri-

cally conductive fillers on PTC effects is presented elsewhere

[65]. PTC effects were observed in polymer composites of both

micro-size fillers, such as CB [65], and nano-size fillers, such

as MWCNT [66] and were attributed to thermal expansion of

the matrix polymer [67] and the mobility of the polymer

chains [68]. Although PTC effect was initially observed in

semi-crystalline polymers [64,65], cross-linked, amorphous

polymer networks also showed PTC effect [69]. A very weak

PTC effect was observed for linear amorphous polymers [70].

A number of researchers observed strong correlation be-

tween crystallinity and PTC effects. Meyer [68] showed that

PTC behavior was possible only with semi-crystalline poly-

mers; uncrosslinked amorphous polymers did not exhibit

PTC effects. The PTC effect is diminished in composites of

semi-crystalline polymers if specimens are quenched rapidly

to yield low crystallinity [71] or if crystallinity of specimens is

reduced by thermal degradation [72]. On the other hand, in-

creased crystallinity of matrix polymer due to annealing pro-

motes PTC effects [73].

In this work, the PTC effects in SMPU composites were

determined from the temperature dependent volume resistiv-

ity data. For this purpose, composites of 3 wt% CNF prepared

in chaotic mixer, 3 wt% CB composite prepared in internal

mixer, and 5 wt% ox-CNF prepared in chaotic mixer were con-

sidered. These composites showed room temperature electri-

cal conductivity r � 10�6–10�5 S/cm (see Fig. 10) and it was

anticipated that the conductive networks in these composites

would be destroyed with melting of the soft segment crystals

and hence resistivity would increase with an increase of tem-

perature. The resistivity versus temperature data presented in

Fig. 12 shows that resistivity of CB composites first increased

up to approximately 45 �C and then reduced. On the other

hand, composites of CNF and ox-CNF did not show PTC effect.

Composites with lower or higher filler contents than the ones

reported in Fig. 12 did not show any significant change in

resistivity upon heating. Similar observation was made by

other researchers [74]. The data reported in Fig. 12 were repro-

ducible during cyclic testing over a temperature interval of

25–60 �C. In order to interpret such trends, it is noted that

the soft segment crystallinity of these composites were 7.6,

12.7, and 16.7 wt%, respectively for CNF, ox-CNF, and CB

Fig. 12 – Temperature dependent volume resistivity of

SMPU/3 wt% CNF, SMPU/3 wt% CB, and SMPU/ox-CNF 5 wt%

composites. The composites prepared in Brabender internal

mixer and in chaotic mixer are represented, respectively, by

(B) and (C).

Fig. 14 – Schematic of crystal melting and the origin of the

PTC effect in the composites of semi-crystalline polymers.

Tmelt represents the melting temperature of the crystals.

Melting of large amounts of crystals augments the volume

and disturbs the connectivity of the conductive particles.

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 991

composites. Also recall that the SMPU composites used in this

work contained 33 wt% hard segments and 67 wt% soft

segments.

Fig. 12 reveals a clear PTC effect in the case of CB/SMPU

composite. The volume resistivity increased by 3 orders of

magnitude as the temperature was increased from 30 to

45 �C. The DSC data on soft segment crystal melting range

(Fig. 13) indicated that PTC behavior occurred in the proximity

of soft segment crystal melting. In view of these, the PTC ef-

fects in CB composites can be explained using a schematic

in Fig. 14. Note that the illustration in Fig. 14 is a simple,

two-dimensional visualization of soft segment crystal melt-

endo

Fig. 13 – DSC trace for CB/SMPU composite with 3 wt% filler

content. Note the melting interval of soft segment crystals

between 30 and 55 �C.

ing and reorganization of conductive networks. It is hypothe-

sized the increase of volume of the matrix due to crystal

melting disturbed the connectivity of CB particles and turned

the composite into an insulator. Pronounced PTC effects were

absent in SMPU composites of CNF and ox-CNF (Fig. 12). Two

possible scenarios might be constructed on why significant

PTC effects were absent in composites of CNF and ox-CNF.

First, the absence of PTC effect in composites of CNF and

ox-CNF can be attributed to lower soft segment crystallinity,

e.g. 7.6% and 12.7% respectively, compared to 16.7% for com-

posite of CB. This hypothesis (scenario I) is schematically pre-

sented in Fig. 15a. Melting of small amounts of soft segment

crystals might have caused a relatively insignificant change

in morphology and hence a negligible disturbance to the con-

nective networks which were present in specimens at room

temperature. Consequently, the conductive filler networks

were preserved and no change in conductivity was observed.

In scenario II, the conductive filler networks in the compos-

ites of CNF and ox-CNF were insensitive to crystal melting

due to the presence of stable, interconnected networks

formed by a large number of nano size, large aspect ratio fill-

ers. This is schematically illustrated in Fig. 15b.

In order to gain further insight on the origin of the PTC

behavior and to evaluate the two scenarios presented above,

we turned our attention to a set of composites of PCL diol.

We determined the temperature dependent volume resistivity

data of CNF/PCL diol and CB/PCL diol composites, each with

3 wt% filler content, and those of ox-CNF/PCL diol composites

with 5 wt% filler content. These data are presented in Fig. 16.

It is seen that all PCL diol composites were highly crystalline

(about 58%, Table 2) and all exhibited clear PTC effects, inde-

pendent of the filler type. The crystallinity of PCL diol com-

posites was much higher compared to those of SMPU

composites. The most striking case was the CNF composites,

whereas the crystallinity of CNF/PCL diol composites was

about an order of magnitude higher than that of CNF/SMPU

Fig. 15 – (a) Scenario I on the origin of the absence of PTC

effects in the composites of CNF and ox-CNF. Tmelt

represents the melting temperature of the crystals. Melting

of small amounts of crystals might have caused a relatively

insignificant increase in volume and hence has limited

effect on the connectivity of the network. (b) Scenario II on

the origin of the absence of PTC effects in the composites of

CNF/SMPU and ox-CNF. Tmelt represents the melting

temperature of the crystals. Due to high aspect ratio and

large number of particles per unit volume, crystal melting

might have a relatively insignificant impact on the

connectivity of the network.

Ω

Fig. 16 – Temperature dependent volume resistivity PCL diol

composites. CNF/PCL diol and CB/PCL diol composites have

each 3 wt% filler content, and ox-CNF/PCL diol composite

has 5 wt% filler content.

Fig. 17 – Typical change in specimen length due to thermal

expansion behavior of CB/SMPU composite with 3 wt% filler

content.

992 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

composites. The data on crystallinity and PTC effects in SMPU

and PCL diol composites suggested that the PTC effects are

associated with the melting of relatively large amounts of

crystals (scenario I). Proposing a connection between PTC ef-

fects and the level of crystallinity in the matrix is also in

accordance with the previous observations reported in litera-

ture whereby strong PTC effects were observed in semi-crys-

talline polymers, however not in amorphous polymers [64,70].

3.5. Relationship between thermal expansion and PTCeffect

Some investigators attribute the PTC effects in semi-crystal-

line and amorphous polymers to thermal expansion of the

matrix polymer during heating [67]. In view of this, thermal

expansion behavior of CB/SMPU composite was investigated

to determine its relationship to the trend observed in Fig. 12.

The change in sample length with temperature for CB/

SMPU composite with 3 wt% filler content is presented in

Fig. 17. Three regions are identified in Fig. 17. Region 1, be-

tween about �40 and 0 �C, covers the temperatures below

the melting range of the soft segment crystals. Region 2 de-

fines a transition regime where the melting of the soft seg-

ment crystals took place between about 0 and 50 �C. Region

3, on the other hand, covers the temperature range higher

than the melting range of the soft segment crystals between

about 50 and 120 �C. The coefficient of thermal expansion

Fig. 18 – The variation of coefficient of thermal expansion

(CTE) with temperature in CB filled SMPU composite with

3 wt% filler. CTE in regions of linear thermal expansion are

designated as coefficient of linear thermal expansion (CLTE).

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 993

(Fig. 18) was determined from the slope of the sample length

versus temperature curves. A reflection into Fig. 18 revealed

that thermal expansion was linear in Regions of 1 and 3,

while the specimen dimension changed non-linearly with

temperature in Region 2. Consequently, the coefficient of

Fig. 19 – (a) Schematic illustration of CB particles excluded by cr

temperature of the soft segment crystals. Here, do represents th

have a concentration near to the percolation threshold. Ro stands

matrix polymer. (b) Schematic illustration of the excluded CB pa

network is lost due to increased gap between CB particles resul

the morphology just after crystal melting, whereas CB particles

represents the gap between two adjacent CB particles and Rf stan

polymer.

thermal expansion in Regions 1 and 3 is specifically desig-

nated as coefficient of linear thermal expansion (CLTE). A

comparison of the effects of temperature on coefficient of

thermal expansion (Fig. 18), with those on soft segment crys-

tallinity (Fig. 13) and those on volume resistivity (Fig. 12) for

CB/SMPU composite with 3 wt% filler gave more insight on

the origin of the PTC behavior. The data presented in Figs.

12, 13 and 18 indicate that the soft segment crystal melting

is mainly responsible for the non-linear thermal expansion

and also for the PTC effect. Also note that the onset temper-

ature of PTC effect (�30 �C; Fig. 12) was much higher than

the onset of non-linear thermal expansion (�0 �C; Fig. 18). A

similar mismatch of onset temperatures of non-linear ther-

mal expansion and PTC effect was reported by other research-

ers [75]. It was also reported that PTC effect is triggered

beyond a critical thermal expansion level [72].

3.6. Analysis of length scales associated with PTC effects

The local morphology of CB/SMPU composite with 3 wt% filler

content could be considered as the one depicted in Fig. 19a, at

temperatures below the crystal melting temperature of the

soft segments. For the sake of simplicity of analysis, CB parti-

cles are thought to be located on the surface of a spherical do-

main of the semi-crystalline polymer matrix. The radius of

the sphere and the distance between two arbitrary CB parti-

cles are designated by, respectively, Ro and do. Note that the

ystalline domains, at temperatures below the melting

e tunneling gap between two adjacent CB particles which

for the radius of the hypothetical sphere of semi-crystalline

rticles at the time instant when the connectivity of the

ting from crystal melting. The schematic may correspond to

were not redistributed in the amorphous matrix. Here, df

ds for the radius of the hypothetical sphere of molten matrix

994 C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7

distance do is the minimum required gap distance between

two adjacent conductive filler particles necessary for electron

tunneling, since the filler content in the composite is at about

the percolation threshold. Although the schematic presented

in Fig. 19a is a greatly simplified view of the true morphology

of CB composites, it is not unrealistic. The tunneling conduc-

tivity occurs almost exclusively between two neighboring par-

ticles and any contribution from a particle which is not

nearest neighbor is negligible, as illustrated by the strong

exponential dependence of tunneling current on the gap dis-

tance [67]. Fig. 19b illustrates the change in local morphology

depicted in Fig. 19a just at the moment of crystal melting. The

volume of the specimen increases upon crystal melting and

the distance between two adjacent particles increases. At this

point, the connectivity is lost and the composite turns to an

insulator. Note that at longer elapsed time after crystal melt-

ing, the filler particles may be redistributed in the matrix.

Using the geometry considered in Fig. 19a and b and the

CTE values presented in Fig. 18, the relation between do and

df can be approximated as do = 1.004df.

A separate calculation can also be made using the volume

change upon melting. The specific volume of a polymer in

amorphous state is generally about 10% higher than that of

the same polymer in crystalline state [76]. This is also true

for PCL diol, which has an amorphous state density of

1.094 g/cm3 and a crystalline density of 1.187 g/cm3 [77]. Using

these density data and the geometry considered in Fig. 19a

and b, and noting that the soft segments are only partially

crystalline in SMPU, the relation between do and df can be also

determined as do = 1.004df. Hence, the results of two separate

calculations, one with CTE presented in Fig. 18 and, one with

the density values reported in literature coincide closely with

each other. These results support the reasoning that the crys-

tal melting causes the pronounced thermal expansion. Thus,

the origin of PTC effects in our system could not be attributed

only to crystal melting or only to thermal expansion, but a

combination of the two.

Now let us turn our attention to the length scales associ-

ated with PTC effects. The tunneling gap (do) in carbon black

filled elastomers was estimated to be in the range of about 2–

10 nm by different researchers [78–80]. In view of this value

for do and the expression that we obtained for the relation-

ship between do and df, the necessary increase in tunneling

distance to convert the conductor to an insulator should be

on the order of 0.1–0.5 A. Both Voet [65] and Medalia [67] pre-

dicted that only small changes in relative positions of carbon

black particles would result in break down of connectivity and

could lead to strong PTC effects. The calculations of Sherman

et al. [81] also showed that the conductivity is extremely sen-

sitive to the changes in the tunneling distance. Our relatively

simple calculations on the length scale associated with PTC

effects agree with the previous predictions.

3.7. Negative temperature coefficient (NTC) effect incomposites

A few words are in order on the negative temperature coeffi-

cient (NTC) effect that we observed in both SMPU and PCL diol

composites. NTC effect is characterized by the restoration of

electrical conductivity upon further heating of the specimens

[82]. It is thought to be related to filler mobility and filler redis-

tribution with continued heating. Cross-linked polymer net-

works usually do not show NTC due to reduced filler

mobility, although the reasons for this decreased mobility

were not elaborated [65]. It may originate solely from the phys-

ical confinement of filler particles in a thermally stable, cross-

linked network or possibly from the chemical bonding of the

filler particles with the polymer network [83]. Our data in Figs.

12 and 16, respectively, above 52 and 48 �C, indicate that both

SMPU and PCL diol composites showed NTC effect. This was

expected since both materials were comprised of linear ther-

moplastic polymer chains and did not possess any thermally

stable network to hinder mobility of the conductive fillers.

3.8. Changes in electrical resistance induced by tensilestrain

A critical issue in actuating SMP composites by resistive heat-

ing is the stability of electrical conductivity as the materials

undergo large strain – as high as several hundred percents –

in typical shape memory cycles. The effects of uniaxial tensile

strain on electrical conductivity were widely investigated in

the context of CB filled conductive elastomers [84–86]. A sig-

nificant strain induced increase in resistance was observed

in CB filled elastomers, especially at filler concentrations

around the percolation threshold. This increase was attrib-

uted to disruption of conductive filler networks upon applica-

tion of strain [84]. Elastomeric composites of carbon fiber

showed an opposite behavior [87]; in this case, the magnitude

of change in resistivity and the rate of the change showed

weak dependence on strains. This was attributed to weak

interfacial bonding between carbon fiber and the polymer;

the interfacial bonding was much stronger in CB filled com-

posites [87].

The resistance of composites prepared in chaotic mixer

and in Brabender internal mixer as function of uniaxial ten-

sile strain is presented in Fig. 20. The measurements were

performed at 60 �C, the stretching temperature of specimens.

In addition, the soft segment crystals were all melted at 60 �C.

The composites with filler content approximately at the per-

colation threshold were considered in Fig. 20. Composites

with lower or higher filler contents did not show any signifi-

cant change in resistance upon tensile deformation, as re-

ported previously by other researchers [86]. It was observed

that CB/SMPU composites with 5 wt% or higher CB content

failed beyond 10% elongation and, therefore was not consid-

ered. It is noted from Fig. 20 that the resistance in both axial

and transverse directions of CB/SMPU composite containing

3 wt% CB increased linearly with strain up to 50% strain,

although the changes in resistance in axial and transverse

directions were different (Fig. 20); similar observation were re-

ported by other researchers [88]. The difference between

resistance in axial and transverse directions was attributed

to preferred orientation of carbon black particles in the strain

direction [84]. However, the resistance of CNF/SMPU and ox-

CNF/SMPU composites did not change significantly with

strain. Paik et al. [8] also reported marginal changes in resis-

tance of MWCNT/SMPU composite with 3 wt% filler content.

Several factors contribute to such differences in resistance

vs. strain behavior of composites of CNF, ox-CNF, and CB.

1.00E+00

1.00E+02

1.00E+04

1.00E+06

Fig. 20 – The ratio of resistance R/Ro of SMPU composites as

function of tensile strain, where Ro is the original resistance

and R is the resistance of stretched specimen. The

composites prepared in Brabender internal mixer and in

chaotic mixer are represented, respectively, by (B) and (C).

The filler content in the composites is given as weight

percent.

C A R B O N 4 7 ( 2 0 0 9 ) 9 8 1 – 9 9 7 995

First, this can be attributed to differences in interfacial

strengths between CNF, ox-CNF, and CB and the SMPU as sug-

gested previously [87]. Second, high aspect ratio fillers such as

CNF and ox-CNF possibly maintained enough connectivity so

as to preserve the conductive networks even at large strain.

Third, CB particles of microscopic size presented much less

filler surface area per unit volume in the composite compared

to nanofillers, such as CNF and ox-CNF. A conductive network

formed by a larger number of individual filler particles might

be more stable under strain compared to a network formed by

relatively few, but larger particles.

Recall that the data presented in Fig. 20 were obtained at

60 �C, whereby the soft segment crystals were all melted. Also

recall that specimens were cooled down to room temperature

in a part of the shape memory testing cycle. It became appar-

ent that some of the conductive networks destroyed at the

time of tensile elongation might reform due to crystallization

of the soft segment crystals [63] and the composite would ex-

hibit electrical conductivity. In view of this, the resistance of

stretched specimens was measured during cooling to room

temperature. The cooling rate was about 15 �C/min as dis-

cussed elsewhere [12,45]. However, it was observed that the

resistance of the composites did not show appreciable

changes upon cooling. In the case of composites of CNF and

ox-CNF, the growth of soft segment crystals did not affect

the resistance. Data also suggested that in the case of CB

composites, the connectivity of filler network was already lost

due to tensile elongation and crystallization was not able to

restore the conductive networks of CB particles.

4. Conclusions

The following conclusion can be drawn from this study. The

soft segment crystallinity of SMPU reduced in the presence

of CNF and ox-CNF. The reduction was more severe in the

presence of CNF due to non-polar nature of CNF in a polar

SMPU. The reduction in the soft segment crystallinity was

responsible for the absence of PTC effects in composites of

CNF and ox-CNF, although composites of CB with almost

the same filler loading showed pronounced PTC effects. This

observation suggested that a minimum level of matrix crys-

tallinity is required for shape memory effects. A close rela-

tionship between abrupt thermal expansion of SMPU

matrix and PTC effects was determined. It was also sug-

gested that the length scales associated with PTC effects

were significantly small, on the order of 0.1–0.5 A. It was also

observed that CNF/SMPU and ox-CNF/SMPU composites had

highly stable electrical conductivity under uniaxial tensile

strain, as compared to CB/SMPU composites. This property

of CNF and ox-CNF composites has central importance for

shape memory applications, whereby a stable conductivity

under strain is a must for actuation by resistive heating. A

corollary of this study was that a high level of crystallinity

may cause a PTC effect and prevent any actuation through

resistive heating. However, a carefully tailored compound

which has reduced crystallinity and which requires mini-

mum amount of filler may prevent PTC phenomenon and

could supply necessary electrical conductivity over the oper-

ating temperature range, while offering enough soft segment

crystallinity and rubberlike properties for excellent shape

memory function. This study also established the impor-

tance of chaotic mixers as a powerful tool in preparation

of SMPU and SMPU composites. Better mixing under rela-

tively low-shear rates in chaotic mixers supplied SMPU

nanocomposites of CNF and ox-CNF with higher soft seg-

ment crystallinity and higher electrical conductivity with

lower percolation thresholds.

Acknowledgements

Partial financial support for this work was provided by Na-

tional Science Foundation in the form of CAREER Award

(DMI-0134106) to S.C.J. Authors gratefully acknowledge Dr.

Rex Ramsier and Justin Walker for their help with XPS data

and Ms. Nergis Makal for her help with NMR data.

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