The mechanical behaviour of ZnO nano-particle modified

styrene acrylonitrile copolymers

K. Zuoa, B.R.K. Blackmana,*, J.G. Williamsa,c, H. Steiningerb

a Imperial College London, Department of Mechanical Engineering, City & Guilds Building,

South Kensington Campus, London SW7 2AZ, UK

bBASF SE, Materials Physics and Analytics, 67056 Ludwigshafen, Germany

cAMME Department, University of Sydney, Australia

Keywords: A. Nano composites; B. Fracture toughness; C. Crack

Abstract

Two amorphous SAN copolymers with acrylonitrile contents of 24% and 34% have been

modified by the incorporation of ZnO nano-particles at volume fractions of up to 2.00 %. For

the SAN with 24% acrylonitrile content, three types of nano-particles were added. The first two

were cylindrical (nano-rods) and the third were spherical. For the SAN with 34% acrylonitrile

content, just one of the cylindrical nano-particle types were used. The dispersion of the nano-

particles was assessed using atomic force microscopy and agglomeration was observed for the

higher volume fractions. Young's modulus and fracture toughness values were measured for

each system as a function of volume fraction. The smaller nano-rods were found to be the most

effective for both stiffening and toughening the polymers and the spherical particles were found

not to contribute to stiffness due to poor particle-matrix adhesion. At room temperature the

maximum toughening contributions for all particles occurred at volume fractions of about 0.1%,

above which a reduction was observed. At 80 °C the toughening contributions continued to rise

slowly past 0.1% volume fraction. The toughening is consistent with the notion that the

particles debond from the matrix and then plastic hole growth around the detached particles is

initiated. The peak toughening observed at 0.1% volume fractions are the result of

agglomeration occurring at higher volume fractions.

1. Introduction

1

In recent years the modification of polymer properties via the incorporation of nano-particles

has been a very popular research topic. Much of the literature has focused on the addition of

spherical nano-particles to thermosetting matrices [1-3], where relatively large quantities of

particles can be well dispersed, without agglomeration occurring. Typically it is more difficult

to disperse nano-particles into thermoplastic matrices. Extrusion mixers and particle surface

treatments are widely used techniques to facilitate this process [4]. However, agglomeration of

nano-particles remains a problem in thermoplastics at relatively low particle volume fractions

and thus efforts continue in the field to improve this situation. The use of different particle

shapes and aspect ratios can have an effect on dispersion, for example spherical particles may be

replaced with cylindrical, rod-like particles of various diameters and aspect ratios, and there has

been much interest also in the use of carbon nano-tubes [5].

Once the nano-particles are dispersed into the polymer matrix, many researchers have reported

that significant stiffening and toughening can occur. The stiffening requires load transfer to take

place between the relatively low Young’s modulus polymer matrix and the usually much higher

Young’s modulus particles. The simple rule of mixtures can then be used to predict the

resulting Young’s modulus of the nano-composite. Although there are complications (e.g.

aspect ratio and orientation effects and the potential agglomeration of particles) the basic

mechanism for stiffening is well understood. However, a universally accepted explanation for

the improved toughness values is currently lacking. Workers have identified a number of

mechanisms that contribute to the toughness of nano-composites and efforts have been made to

determine their individual contributions and also their combined effects. The toughening

mechanisms that have been proposed include particle-matrix debonding, the subsequent plastic

void growth [6], shear yielding (or shear banding) of the matrix between the particles [1, 7], and

in some cases crack tip bridging [5] (e.g. when nano-tubes are employed). Also, more recently

workers have employed a multiscale modelling strategy to describe the toughening of polymers

by nanofillers with different properties being ascribed to the matrix and interphase regions

around particles [8-10]. Such work has considered the combined effects of debonding, plastic

yielding of nano-voids and shear banding in the polymer matrix [10]. The model proposed in

[6] to describe debonding followed by plastic void growth has been used to determine the

energy density of matrix shells around debonding particles using a geometrical description of

particle-particle interactions [11, 12].

In the present work, two thermoplastic copolymers SAN with 24% acrylonitrile and SAN with

34% acrylonitrile were modified with three different nano-particle types - two rod-like,

2

cylindrical particle types and spherical particles. The enhancements in stiffness and toughness

have been measured experimentally in a series of tests at 23 °C and 80 °C.

The model proposed in [6] to describe nano-composite toughening enhancement due to the

addition of spherical particles by (i) particle-matrix debonding followed by (ii) plastic void

growth was extended to accommodate toughening with nano-rods and nano-tubes [13]. Here

experimental data for SAN toughened with nano-spheres is analysed using the spherical model

[6], and also SAN toughened with two sizes of nano-rods are analysed using the analogous

model for rod-like particles [13]. The models define a toughness enhancement factor, X, which

is determined from experiment (for either spheres or rods). The measured values of X (up to the

limiting volume fraction when agglomeration occurs) have been used to determine the model

parameters Ga (the energy for debonding) and the critical interfacial stress, σ c, which controls

the subsequent plastic void growth from the debonded particle.

2. Materials

Two grades of amorphous SAN copolymers were used in this study. The first had a trade name

of Luran VLN and contained 24% acrylonitrile in the copolymer. This is referred to as SAN-24

in the present work. The second had a trade name of Luran VLP and contained 34% acrylonitrile

in the copolymer. This is referred to as SAN-34 in the present work.

Three types of ZnO nano-particles were employed as fillers, denoted here as Types A, B and C.

Types A and B were nano-rods and Type C were nano-spheres. Type-A nano-rods had an

average diameter of 40 nm with an aspect ratio (length/diameter) of 4, Type-B nano-rods had an

average particle diameter of 12 nm with an aspect ratio of 3, and Type-C particles were

spherical with an average diameter of 10 nm. All materials were supplied by BASF SE,

Germany. The dimensions of the particles are summarised in Table 1.

The particles were surface modified with trioxadecanoic acid (TODS) or with one of two

different types of silane before blending in order to improve the dispersion of the particles into

the matrix. The SAN and ZnO blends were compounded in a Werner & Pfleiderer ZSK-30 twin

screw extruder and the melt was extruded through a flat sheet die. Prior to compression

moulding, the raw sheets were dried at 80 °C in vacuum for 36 hours. The resulting sheets were

then compression moulded into 6 mm and 4 mm thick plates using a hot compression moulding

3

machine (DAKE, USA) at 200 °C and 5 MPa. The composite systems studied are listed in Table

2.

Concerning the nomenclature used in the present paper, when nano-particles are added to the

system, the volume fraction of particles is shown in the system designation preceding the

particle type, e.g. system "SAN-24-0.5A" contains 0.5% by volume fraction of Type-A

particles. As will be discussed in the next section, for systems modified with the Type-B

particles, tests were performed at two temperatures; namely 23 °C and 80 ºC.

3. Experimental

3.1 Thermal analysis

The thermal behaviour of the unmodified and particle modified composites was studied via

dynamic mechanical thermal analysis (DMTA). Tests were performed on bars of dimensions 50

× 3 × 2 mm in a ‘PerkinElmer DMA 8000’ apparatus. Dual cantilever geometry was employed,

at two different frequencies of 1 Hz and 10 Hz over a temperature range from 20 °C to 140 °C.

At least two specimens were tested for each formulation, and the was defined to be the

temperature at the peak value of the loss factor, , as specified in the ISO standard [14].

3.2 Microscopy

To investigate the particle dispersion in the various systems, atomic force microscopy (AFM)

analysis was undertaken using a ‘MultiMode’ scanning probe microscope from Veeco (Santa

Barbara, USA) equipped with a ‘J’ scanner and a ‘NanoScope IV’ controller. A smooth surface

was first prepared by cutting the samples on a ‘cryo-ultramicrotome’ at 23 °C. The scans were

performed in tapping mode using etched silicon probes supplied by Veeco and both height and

phase images were recorded simultaneously.

To allow an investigation of the fracture surfaces following fracture testing, a ‘Hitachi S-3400N

VP’ scanning electron microscope was used. Typically, an accelerating voltage of 15 kV was

used, and the working distance was approximately 10 mm. The fracture surfaces were sputter

coated with a very thin layer of gold (about 15 nm thick) to reduce charging effects prior to the

SEM examinations. The SEM images could show the morphology of the specimens, and also

the roughness of the fracture surface and the fracture path.

4

3.3 Tension, compression and fracture testing

Specimens for testing were prepared in both the longitudinal and transverse directions on the

mouldings with respect to the previous direction of extrusion and no effects of orientation were

observed in the results. Tensile tests were performed on an Instron testing machine (model

5584) running Bluehill™ control software. The tests were conducted in accordance to ISO 527

[15] standards to determine the Young’s modulus, E. Dumbbell samples, approximately 4 mm

wide × 4 mm thick were used. At least five specimens were tested for each formulation. All tests

were performed using a contact extensometer with a gauge length of 25 mm, at a crosshead

speed of 1 mm/min.

Uniaxial compression tests were performed to determine the yield stress, σ Y, since testing of

relatively brittle materials such as SAN in tension can lead to brittle failure prior to the yield

point. During uniaxial compression, brittle fracture was avoided and plastic yielding was

observed prior to failure. The compression specimens were cylindrical, 6 mm in height (parallel

to the direction applied force), and 20 mm in diameter. A synthetic grease was used to reduce

friction between the specimens and the compression plates. The specimens were tested at a

constant crosshead speed of 1 mm/min according to ISO 604 [16].

Fracture tests were performed on single edge notch bending (SENB) specimens according to

ISO 13586 [17] to obtain the mode I fracture toughness, , and the fracture energy, , at

fracture initiation. Natural pre-cracks were produced by tapping new razor blades into the notch

tips prior to testing. The SENB specimens were tested at a constant displacement rate of

1 mm/min, and at a test temperature of 23 ºC or 80 ºC. Indentation correction tests were

performed as specified in [17].

Tensile, compression and fracture tests were also conducted at 80 °C. All tests at 80 °C were

carried out in an environmental chamber. The test chamber allowed for an accurate temperature

control (±1 K) by using two thermocouples; one thermocouple was inserted into a reference

sample, while the other was freely suspended in the chamber. The chamber was heated up

gradually to 80 °C, and all specimens were stored in the heated chamber for 1 hour prior to

testing to reach thermal equilibrium. Once a specimen was mounted and the chamber door

closed, a soaking time of at least 10 minutes elapsed, ensuring the temperature in the chamber

had stabilised.

5

In order for the fracture results to be considered valid [17], it is required that the size criteria for

plane strain conditions be satisfied:

B ,a , (W −a )>2.5( KQ

σY)

2

(1)

where B ,W and aare the thickness, width and crack length (pre-notch plus razor crack) of the

specimen, respectively. KQ is the conditional or trial value of K c and σ Y is the yield stress of the

material at the appropriate temperature and loading rate.

4. Results

4.1 Particle dispersion

The morphology study was undertaken by using atomic force microscopy (AFM). Fig. 1 shows

AFM phase images of SAN-24-A, SAN-24-B and SAN-24-C composites at various particle

volume fractions. The nano-particles are seen as ‘lighter regions’ on the phase image as the

stiffness of ZnO particles is significantly greater than that of the surrounding polymer. The

microstructure of unmodified SAN was featureless at this scale. The Type-A nano-rods were

difficult to disperse into the matrix. With the addition of 2.00 vol.% of Type-A nano-rods,

agglomerates as large as 300 nm were observed in the composite (Fig. 1 (b)). The Type-B nano-

rods and Type-C nano-spheres were relatively well dispersed through the matrix up to 0.30 vol.

%. Fig. 1 (c) & (e) show the good dispersion achieved for 0.10 vol.% Type-B and Type-C

particles in SAN-24. However, agglomerates as large as 80 nm and 120 nm were observed in

composites containing 0.50 vol.% Type-B nano-rods (Fig. 1 (d)) and 1.25 vol.% Type-C nano-

spheres (Fig. 1 (f)), respectively. Similar observations were made for the system comprising

the SAN-34 matrix with the Type-B nano-rods (note systems SAN-34-A and SAN-34-C were

not tested).

4.2 Thermal behaviour

The storage modulus, , loss factor, , and glass transition temperature, , of the

unmodified and particle modified systems determined by using DMTA are shown in Fig. 2.

6

Typical traces for the SAN-24-0.5B system tested at 1 Hz and 10 Hz are shown. The value of

measured was frequency dependent reflecting common viscoelastic materials response,

which was also observed by Johnsen et al. [18] and Brooker [14] . The obtained at 10 Hz

was about 6 K higher than that obtained at 1 Hz. The of unmodified SAN was measured to

be 116 ± 2 °C at 1 Hz, and 122 ± 2 °C at 10 Hz. The addition of ZnO nano-particles did not

significantly affect the glass transition temperatures of the modified SAN composites, as is

shown in Table 3.

4.3 Tensile testing results

For the tensile tests carried out at 23 °C, no obvious plastic deformation was observed for the

unmodified and nano-modified SAN composites and all materials fractured in the elastic region.

The values of the tensile fracture stress, σ f , are given in Table 3 and show significant decreases

with volume fraction for the A and C particles. However, the values remain approximately

constant for the B particles. The addition of nano-particles is expected to stiffen the composite

due to the high modulus of the particles (approximately 140 GPa [20]). An average tensile

modulus of 3.87 ± 0.05 GPa was measured for the unmodified SAN-24 at 23 °C, in good

agreement with that reported by Cebon et al. [15] for this thermoplastic. Fig. 3 plots the values,

for the various modified SAN-24 composites, of normalised Young’s modulus,( Ec

Em−1), against

the volume fraction of nano-particles added, where Ec and Em are the Young’s moduli of the

modified composite and the control matrix respectively. The results are presented in this form so

that the enhancement in the property, i.e. Young’s modulus or fracture toughness below, is

emphasised. The lines drawn are the best linear fits to the experimental values. It can be seen

that the Type-B nano-rods were most effective over the range of volume fractions studied,

giving the greatest slope on Fig 3. This will be discussed further in Section 5. Interestingly, the

moduli for the SAN-24-C system fall somewhat below the unmodified control. It is thought that

this is due to the nano-particles being only very weakly bonded to the matrix in this system, and

hence failing to support the particle-matrix load transfer.

7

4.4 Fracture testing results

In all the tests the cracks initiated suddenly from the pre-crack, failure thus occurred in a brittle

manner. The values of K c and Gc were obtained at fracture initiation as defined in ISO 13586.

All tests passed the linearity and plane strain validity crosschecks (Eq. 1). The Gc values were

calculated using the energy method as described in [17].

The fracture results at 23 °C are presented in Table 3. The normalised values of fracture energy

( Gc

Gm−1), where Gc and Gm are the fracture energies of the composite and the matrix

respectively were plotted against the nano-particle volume fraction in Fig. 4 for the SAN-24

matrix. Although there is a limited amount of data at very low volume fractions, it would appear

that Type B nano-rods were the most efficient of the three particle types in toughening the

matrix. Also for each of the three particle types, the maximum toughness was achieved at a

particle volume fraction of about 0.10 %. For each particle type the fracture toughness then

decayed very quickly with increasing volume fraction after the maximum toughness was

achieved. Type-C nano-spheres produced similar toughening at the maximum to Type-A nano-

rods, but the values then decayed more rapidly post maximum. The dramatic decrease in

toughness following the maxima was caused by the formation of nano-particle agglomerates at

the higher volume fractions, as depicted in Fig. 1.

The results for the two different copolymers at 80 ˚C, are summarised in Table 4. The

copolymer with higher acrylonitrile content (SAN-34) possessed the higher toughness. The

values of Gc for the unmodified control specimens for SAN-24-0 were 217 J/m2 compared to

561 J/m2 for the SAN-34-0 control specimens. Thus, the SAN-34 matrix was significantly

tougher the SAN-24 matrix. However, on a normalised basis, the trends in toughness

enhancement were similar in the SAN-24 and SAN-34 matrices when modified with the Type-B

nano-rods although proportionately the greater enhancement was seen in the more brittle system.

The results are shown in Figure 5.

Comparing the results for the SAN-24 and SAN-34 matrices at 80 ˚C, then for both matrices the

higher temperature led to a change in the toughening trends, with the toughness of the systems

now continuing to rise, although slowly, with increasing volume fraction. The results for SAN-

24-B are shown in Figure 6(a) and the results for SAN-34-B are shown in Figure 6(b).

8

The micrographs obtained in the scanning electron microscope are shown in Figures 7 and 8 for

the systems modified with the Type-B nano-rods and tested at 23 °C and 80 °C respectively.

Figure 7(a) shows the featureless fracture surface from the SAN-24-0 control tests and Figure

7(b) shows the fracture surface following testing of the SAN-24-0.05B specimens. Circled in

this figure are sites where evidence exists for the cavitation and then subsequent fibrillation of

the SAN-24 matrix following addition of Type-B nano-rods, i.e. a crazing mechanism. Figures

7(c) and (d) indicate the presence of conic markings on the surfaces of the SAN-34 specimens

and as these are also present in the unmodified control, they are not due to the presence of the

Type-B nano-rods. Figure 7(d) clearly shows the presence of voids on the fracture surface

following the addition of the Type-B nano-rods. These voids were measured to be

around 100 nm to 200 nm in diameter, i.e. much larger than the Type-B nano-rods used to

reinforce the matrices. These voids have most likely resulted from the debonding of the particle-

matrix interface and the subsequent plastic hole growth. Figure 8 shows the resulting fracture

surfaces for SAN-24 and SAN-34 matrices reinforced with Type-B nano-rods at 80 °C. With

the exception of the SAN-24 control specimen, all surfaces showed the presence of voids with

sizes in the range of 100 nm to 500 nm. Some of the voids are indicated by arrows in the figure.

Again, this would suggest that debonding of the particle-matrix interface and then subsequent

plastic hole growth had occurred. This has been observed by several other workers and will be

explored quantitatively in the next section.

5. Discussion

5.1 Tensile modulus results

If particles are well bonded to the matrix then we would expect a simple law of mixtures for low

volume fractions (ϕ≪1):

Ec ≈ Em+Ep ϕ

where Em is the modulus of the matrix i.e. approximately 4 GPa and Ep is the modulus of the

particle i.e. approximately 140 GPa, so that:

( Ec

Em−1)≈ 35 ϕ

This line is drawn in Fig. 3 and shows that at very small ϕ the Type-B particles follow this line,

implying the particles are adhered. However, for the larger volume fractions, the slope decreases

9

and falls to a value of 11.6 when all the points are included. This would be consistent with a

reduction in the particle to matrix adhesion leading to the subsequent significant deviation from

the rule of mixtures for all but the smallest value of . The Type-A particles followed a line

with a slope of 5.2 and the Type-C particles indicate no adhesion (in fact, there is negative

enhancement with Type-C particles, an observation which is repeated for the toughness

measurement and is discussed further in the next section. The slopes suggest that the Type-B

particles were well adhered for low ϕ but at higher values the data suggest that about 2/3 were

agglomerated and were thus not adhered as the slope fell by this proportion. For the Type-A

particles a similar argument suggests that the fraction of particles agglomerated is about 0.84

and for Type-C particles all are agglomerated.

5.2 Fracture toughness results

The enhancement in toughness resulting from the addition of the various nano-particles to the

SAN copolymers is considered to result largely from the plastic void growth which follows the

debonding of the particle from the surrounding matrix. This mechanism was shown to be

important for the toughening of epoxy with ATH and silica particles and also for the toughening

of polypropylene with CaCO3 [21].

The mechanism was modelled [6] as a two stage process by firstly considering rigid spherical

particles bonded to the matrix. On loading the specimen the particles in the crack tip region

undergo an increase in interfacial stress until the particles debond. The number of particles

participating in this process depends upon the particle volume fraction and the size of the crack

tip process zone. This debonding process dissipates the interfacial energy per unit area of

particle, Ga. Although this is a small amount of the total energy dissipated, the debonding step is

an essential prerequisite for stage two- given by the resulting void in the matrix expanding.

This process dissipates plastic energy. The two energy dissipating processes together cause the

toughness of the composite to increase from the value of the matrix (when no particles are

present, i.e. ϕ=0) to Gc (when the particles are present at a volume fraction of ϕ. The

enhancement in toughness described by the model is given by:

( Gc

Gm−1)=Xϕ (2)

10

where the toughening factor, X , is a function of Ga, the particle radius ro, the matrix yield stress

σ Y and Young’s modulus of the matrix Em. This analysis has been extended [13] to rigid, rod-

like particles where the same two stage mechanism pertains. In this case, the appropriate

relationships for rod-like particles with an aspect ratio of 3, as used here for the Type-B

particles, and for ν=1/3 , are given by:

X=0.584( ex−1

x−0.318) (3)

where

x2=3.25( Em Ga

σY2 r o

)=( 32

σ c

σY)

2

(4)

The procedure followed is that X is determined from the experimentally measured values of Gm

and Gc according to eqn. (2). Eqn (3) is then used to determine the critical stress ratio factor, x,

and eqn (4) is then used to determine the interfacial energy, Ga and also the critical stress for

debonding, σ c. Therefore, if the particle size, matrix properties and the interfacial energy are all

known, the model can be used to predict the toughness enhancement of the composite. In the

present work, the toughness enhancements have been determined experimentally, so the model

has been used to output values of x ,Ga and σc

σY . For each of the materials investigated, the

maximum toughness enhancement occurred at about ϕ=0.1 % (Fig. 4) and therefore X has been

determined at this value.

Table 5 gives the values of E, σ Y , ro, the test temperature T, X , x ,Ga and σ c

σY for the SAN-24

and SAN-34 systems. The values of x ,Ga and σc

σY are calculated using eqns. (3) and (4) except

for the Type-C particles which were spherical and used the form appropriate to that case [6].

(Type-A particles had an aspect ratio of 4 but the difference in the equations is negligible).

The x values were determined, being in the range 8.9 to 10 for the three particle types, and the

Ga values were determined to be in the range 0.19 to 1.0 J/m2. The value of Ga for the Type-B

particles at about 0.3 J/m2 are similar to the value obtained for nano-silica particles of the same

size in an epoxy matrix. The value of Ga = 1.0 J/m2 for the composites with Type-A particles is

notably higher than for particles Types B or C. This observation was noted and discussed at

11

length in [6]. It is of interest to note that for spherical particles x=σc

σY whilst for rods x=3

2σc

σY

so that σ c

σY is about 6 for rods and 9 for spheres. This suggests that some of the problem in

toughening with the Type-C particles here (spheres) may arise from these particles requiring

50% higher stress to trigger debonding. For rods the stresses required for debonding are lower

due to their geometry. A ratio of σc

σY≈ 7 was found for silica spheres in epoxy.

The rapid decline in toughness (post maximum) with increasing ϕ may also be explained by this

analysis. It was noted that agglomeration occurred for the higher ϕ values leading to ‘particles’

forming with a radius of up to 60 nm, i.e. effectively increasing the particle size by a factor of

10. Assuming all other parameters remain constant, this decreases the critical stress ratio factor,

x by 10−1 /2=0.32 , i.e. x=3 giving X ≈ 1.2 from eqn. (3). Thus the toughening from the

agglomerates is negligible and the loss of the agglomerated particles from the toughening

process leads to the decline in toughening. The modulus changes also suggest that agglomerated

particles are not adhered, thus preventing toughening.

The analysis of plastic void growth [13] which is controlled by the parameter x also gives an

estimate of the cavity radius, r v, given by:

( rv

ro)

2

≈ 1+32

σY

E c⋅ex−1

Using typical values of x=9 and σY

E≈ 0.02 from Table 5 we have

rv

ro=9.5 . Thus for ro=6 nm,

the average void diameter would be 110 nm. Fig 8. gives a range of sizes from

100 nm – 500 nm and given that the particle size varies, and given also the existence of

agglomerates, the prediction provides some support for the analysis.

An extraordinary feature in Fig. 4 is that for the Type-A and Type-C particles, the values of Gc

Gm

decrease significantly below unity for volume fractions, ϕ>0.5 % . The toughening model

discussed here does not predict such an effect and its source must be sought elsewhere. Some

indication is provided in Table 3 in that the values of tensile fracture stress also decreased with

increasing ϕ which implies a decrease in the crazing stress. It is suggested that the decrease in

Gc arises from the particles disrupting the energy absorbing crazing process in these cases.

12

Finally, it was noteworthy that increasing the test temperature from 23 °C to 80 °C changed the

toughening behaviour of SAN-24 and SAN-34 matrices with the Type-B particles significantly

at the higher volume fractions (Figs 6a and 6b). For both composites, the peak in toughness

which occurred at about a value of =0.1% at 23 °C disappeared, and the toughness continued to

rise further, albeit at a decreasing rate, over the whole range of tested. At temperatures closer

to the glass transition temperature the molecular mobility increases and in its wake, matrix

toughness decreases. This would increase the size of the process zone, 2c, formed ahead of the

crack tip. A larger process zone will contain an increased number of ‘active particles’ i.e. ones

which debond and induce matrix void growth as part of the fracture process. In [13] this was

termed N. In addition, the critical stress, c, for plastic void growth around debonded particles

drops significantly, i.e. by a factor of about 1/3. These effects are summarised in Table 6. As a

result, some aggregates so far excluded from the fracture process now are ‘debonded’, giving

way to subsequent plastic void growth. In the case of aggregates this may not be debonding in

the true sense, but may involve internal fragmentation at stresses below c, from which plastic

void growth can initiate. Hence the density of active filler particles, i.e. those taking part in the

fracture process of the composite by debonding (or fragmentation) and subsequent plastic void

growth, is increased at elevated temperatures. This can be expressed by an increase in the

number of particles in the process zone per unit length, i.e. N/l.

6. Conclusions

Two amorphous SAN copolymers with acrylonitrile contents of 24% and 34% were modified by

incorporating ZnO nano-particles at volume fractions of up to 2.00%. SAN with 24%

acrylonitrile was modified with three types of nano-particles termed Type-A (40 nm diameter

nano-rods with aspect ratio of 4), Type-B (12 nm diameter nano-rods with aspect ratio of 3), and

Type-C (10 nm diameter nano-spheres). SAN with 34% acrylonitrile was just modified with the

Type-B nano-rods. For each composite investigated the dispersion of the nano particles was

assessed using atomic force microscopy and the glass transition temperature was measured

using dynamic mechanical analysis. For each composite, tensile tests and quasi-static fracture

toughness tests were performed at 23 °C. For the composites containing the Type-B nano-rods,

tests were additionally performed at 80 °C. The incorporation of nano-particles was shown to

have no effect on the Tg of the resulting composite and atomic force microscopy showed that

there was significant agglomeration in all the systems for particle volume fractions above 0.1%.

13

It was for this reason that all the composites studied here contain no more than 2% by volume of

added particles.

The tensile modulus results were compared with the prediction from a simple rule of mixtures

and it was shown that there was only agreement for the Type-B nano-rods at volume fractions of

up to 0.1%. For higher volume fractions the modulus fell significantly below that predicted and

this is consistent with the observation that the particles agglomerated in this range. The tensile

modulus results for the composites modified with the larger Type-A nano-rods were all

significantly below the values predicted and that is also consistent with these larger particles

agglomerating at all concentrations (even the smallest in this study at 0.1%). The tensile

modulus results for the composites modified with the Type-C nano-spheres were also

significantly below the values predicted and indeed negative enhancement was observed.

The fracture toughness was measured for each composite at 23 °C and its enhancement was

determined as a function of volume fraction. For all composites there was an initial increase in

toughness with a maximum occurring at about 0.1% by volume of added particles, after which

the enhancements were reversed. For the SAN-24 matrix, Type-B nano-rods were again the

most effective and led to the greatest percentage increases in toughness at +41% relative to the

unmodified matrix (X=410). They were twice as effective as the larger Type-A nano-rods

(which gave a maximum increase in toughness of +19%) , and also were more effective than the

Type-C nano-spheres (which gave a maximum increase in toughness of +27%) . The SAN-34

matrix with a GIc of 936 J/m2 was significantly tougher than the SAN-24 matrix at 316 J/m2 in

the unmodified condition and the addition of Type-B nano-rods gave a maximum enhancement

of +21% (X=210) at 23 °C. When the temperature was increased from 23 °C to 80 °C, the

fracture toughness was shown to rise continuously (albeit at a slower rate) past 0.1% volume

fraction for both matrices when reinforced with the Type-B nano-rods. The fracture results were

further interrogated by the use of the by particle cavitation and void growth toughening model

previously proposed. This model showed that nano-rods were more efficient at toughening than

nano-spheres because the rod-like particles debond at a lower stress than the spheres and can

therefore impart toughness more readily to the composite. This can explain the lower toughness

enhancement achieved by adding the Type-C particles compared to the Type-B particles for an

equivalent particle volume fraction. By inputting the experimentally determined toughness

enhancement into the model, together with the other parameters of yield stress, Young’s

modulus, and particle radius, then the critical stress ratio factor for debonding could be

calculated, together with the parameter Ga, describing the adhesion between the particle and the

14

matrix. The values obtained were sensible and added weight to the validity of the model. The

model was then used to predict how the toughness enhancement would be affected by

agglomeration of particles. The microscopy undertaken here suggested that agglomerates of up

to 60 nm in radius were formed and the model confirmed that this would almost entirely remove

the toughness enhancement. Whilst the toughening model does not predict any negative

enhancement (deterioration) it is suggested that this observation is associated with the

breakdown of the crazing mechanisms which give the higher toughness in thermoplastics.

Finally the difference in toughening behaviour for the SAN matrices modified with B-particles

at 23 and 80 °C for volume fractions greater than 0.1% was rationalised in terms of the greater

number of particles becoming active in the process zone at the higher temperatures.

References

1. Hsieh, T.H., Kinloch, A.J., Masania, K., Sohn Lee, J., Taylor, A.C., Sprenger, S., The

toughness of epoxy polymers and fibre composites modified with rubber microparticles

and silica nanoparticles. Journal of materials science, 2010. 45(5): p. 1193-1210.

2. Liang, Y.L. and R.A. Pearson, Toughening mechanisms in epoxy-silica nanocomposites

(ESNs). Polymer, 2009. 50(20): p. 4895-4905.

3. Wetzel, B., Rosso, P., Haupert, F., Friedrich, K., Epoxy nanocomposites - fracture and

toughening mechanisms. Engineering Fracture Mechanics, 2006. 73(16): p. 2375-2398.

4. Tang, E., Cheng, G., Ma, X.,Zhao, Q., Surface modification of zinc oxide nanoparticle

by PMAA and its dispersion in aqueous system. Applied Surface Science, 2006. 252(14):

p. 5227-5232.

5. Hsieh, T.H., A.J. Kinloch, A.C. Taylor, I.A. Kinloch. The effect of carbon nanotubes on

the fracture toughness and fatigue performance of a thermosetting epoxy polymer.

Journal of Materials Science, 2011. 46: p. 7525-7535.

6. Williams, J.G., Particle toughening by plastic void growth. Composites Science and

Technology, 2010. 70: p. 885-891.

7. Huang, Y. and A.J. Kinloch, Modelling of the toughening mechanisms in rubber-

modified epoxy polymers. Journal of materials science, 1992. 27(10): p. 2763-2769.

8. Zappalorto, M., M. Salviato, and M. Quaresimin. A multiscale model to describe

nanocomposite fracture toughness enhancement by the plastic yielding of nanovoids.

Composites Science and Technology 72 (2012): p.1693-1691.

15

9. Salviato, M., M. Zappalorto, and M. Quaresimin. Plastic shear bands and fracture

toughness improvements of nanoparticle filled polymers: A multiscale analytical model.

Composites Part A 48 (2013): p.144-152.

10. Quaresimin, M., M. Salviato, and M. Zappalorto. A multiscale and multimechanism

approach for the fracture toughness assessment of polymer nanocomposites. Composites

Science and Technology 91 (2014): p.16-21.

11. Lauke, B., Contribution of matrix yielding energy to the crack resistance of particle

reinforced composites. Composites Science and Technology 86 (2013): p.135-141.

12. Lauke, B., and S. Y. Fu, Aspects of fracture toughness modelling of particle filled

polymer composites. Composites Part B 45 (2013): p.1569-1574.

13. Williams, J.G., B.R.K. Blackman, H. Steininger, K. Zuo. Toughening by plastic

cavitation around cylindrical particles and fibres. Composites Science and Technology

103 (2014): p.119-126.

14. ISO 6721-11, (2012) Plastics - Determination of dynamic mechanical properties - Part

11: Glass transition temperature. ISO, Geneva.

15. ISO 527-1, (1993) Plastics-Determination of tensile properties-Part 1: General

principles. ISO, Geneva.

16. ISO 604: (2003) Plastics-Determination of compressive properties. ISO, Geneva.

17. ISO 13586, (2000) Plastics- Standard test method for the determination of fracture

toughness (Gc and Kc)- A linear elastic fracture mechanics approach. ISO, Geneva.

18. Johnsen, B.B., Kinloch, A.J., Mohammed, R.D., Taylor, A.C., Sprenger, S., Toughening

mechanisms of nanoparticle-modified epoxy polymers. Polymer, 2007. 48(2): p. 530-

541.

19. Brooker, R.D., The morphology and toughness of thermoplastic-modified epoxy polymer

with carbon nanotubes or silica nanoparticles. Ph.D. thesis, Department of Mechanical

Engineering, Imperial College London, London, 2008.

20. Cebon, D., M.F. Ashby, L. Lee-Shothaman, CES EduPack 2010 Software. 2010.

21. Stapountzi, O.A., M. Charalambides, and J.G. Williams, Micromechanical models for

stiffness prediction of alumma tri hydrate (ATH) reinforced PMMA: effect of volume

fraction and temperature. Composites Science and Technology, 2009, 69(11-12): p.

2015-2013.

16

Table 1. Details of the ZnO nano-particles used.

Particle Designation Shape Diameter (nm) Aspect ratio (l/d)

A Nano-rod 40 4

B Nano-rod 12 3

C Nano-sphere 10 1

Table 2. The nano-composite systems studied and testing temperatures.

↓Matrix; Particles→ A B C"Composite designation" and (test temperatures)

SAN-24 "SAN-24-A"

(23 ºC)

"SAN-24-B"

(23 ºC & 80 ºC)

"SAN-24-C"

(23 ºC)

SAN-34 - "SAN-34-B"

(23 ºC & 80 ºC)

-

Note: (-) indicates not tested.

17

Table 3. Summary of the values of glass transition temperature, Tg, Young’s modulus, E, yield stress, σy, fracture stress, σf, fracture toughness, Kc, and fracture energy, Gc, for the unmodified and nano-modified SAN composites tested at 23 ºC.

System Nano-content

Tg E σy σf Kc Gc

(vol.%) (°C) (GPa) (MPa) (MPa) (MPa·m1/2) (J/m2)

SAN-24-0 0 116 ± 2 3.87 ± 0.05 88 ± 1 53 1.12 ± 0.08 316 ± 12

SAN-24-A 0.10 117 ± 2 3.89 ± 0.04 88 ± 1 41 1.28 ± 0.10 375 ± 25

SAN-24-A 1.04 116 ± 0 4.06 ± 0.10 88 ± 1 35 1.06 ± 0.03 257 ± 9

SAN-24-A 2.00 115 ± 2 4.28 ± 0.14 88 ± 1 34 0.98 ± 0.04 202 ± 17

SAN-24-B 0.05 117 ± 0 3.95 ± 0.06 88 ± 1 52 1.41 ± 0.02 441 ± 12

SAN-24-B 0.10 115 ± 0 3.98 ± 0.03 88 ± 1 53 1.40 ± 0.03 445 ± 27

SAN-24-B 0.30 116 ± 2 4.03 ± 0.04 88 ± 1 53 1.34 ± 0.02 402 ± 14

SAN-24-B 0.50 116 ± 0 4.06 ± 0.05 88 ± 1 50 1.26 ± 0.05 372 ± 9

SAN-24-C 0.10 116 ± 0 3.85 ± 0.03 89 ± 1 48 1.31 ± 0.03 401 ± 17

SAN-24-C 0.27 117 ± 2 3.75 ± 0.04 89 ± 1 42 1.08 ± 0.06 277 ± 11

SAN-24-C 0.63 116 ± 1 3.74 ± 0.07 90 ± 1 40 0.93 ± 0.03 186 ± 8

SAN-24-C 1.25 116 ± 1 3.77 ± 0.02 90 ± 1 41 0.66 ± 0.04 117 ± 14

SAN-34-0 0 105 4.14 ± 0.03 89 ± 1 75 2.11 ± 0.08 963 ± 25

SAN34-B 0.10 105 4.21 ± 0.07 89 ± 1 77 2.30 ± 0.02 1166 ± 45

SAN34-B 0.50 105 4.24 ± 0.06 89 ± 1 78 2.18 ± 0.07 1043 ± 89

SAN34-B 1.00 105 4.33 ± 0.05 89 ± 1 75 1.93 ± 0.12 814 ± 69

18

Table 4 Summary of the values of Young’s modulus, E, yield stress, σy, fracture stress, σf

fracture toughness, Kc, and fracture energy, Gc, for the unmodified and nano-modified SAN-24

and SAN-34 composites tested at 80 °C.

SystemNano-

contentE σy σf Kc Gc

  vol.% (GPa) (MPa) (MPa) (MPa·m1/2) (J/m2)

SAN24-B

0 3.19 ± 0.03 58 ± 1 34 0.89 ± 0.06 217 ± 17

0.05 3.24 ± 0.07 57 ± 1 35 1.07 ± 0.05 307 ± 12

0.10 3.29 ± 0.06 57 ± 1 36 1.09 ± 0.03 315 ± 12

0.30 3.35 ± 0.07 57 ± 1 33 1.12 ± 0.07 345 ± 15

0.50 3.40 ± 0.05 57 ± 1 37 1.14 ± 0.01 346 ± 12

SAN34-B

0 3.43 ± 0.10 59 ± 1 43 1.49 ± 0.04 561 ± 32

0.10 3.47 ± 0.04 58 ± 2 45 1.54 ± 0.05 632 ± 26

0.50 3.52 ± 0.01 59 ± 2 45 1.58 ± 0.08 677 ± 11

1.00 3.57 ± 0.06 59 ± 1 44 1.60 ± 0.03 699 ± 11

Table 5 Summary of the toughening model input parameters and results for the various material

systems investigated where X, x, Ga, c/y denote: toughening factor, critical stress intensity

factor, interfacial energy and normalised critical stress. Following that sequence the model

parameters are calculated from eqns. (2), (3), (4) and again (4). X is determined from the

toughness enhancement at ϕ=0.1 %.

System T

°C

E

(GPa)

Y

(MPa)

ro

(nm)

X x Ga

(J/m2)

σc

σY

SAN-24-A 23 3.9 88 20 190 8.9 1.00 5.9

SAN-24-B 23 3.9 88 6 410 9.9 0.36 6.6

SAN-24-C 23 3.9 88 5 270 9.7 0.33 9.7

SAN-24-B 80 3.3 57 6 450 10.0 0.19 6.7

SAN-34-B 23 4.2 89 6 210 9.0 0.29 6.0

SAN-34-B 80 3.5 59 6 127 8.4 0.13 5.6

19

Table 6 Comparison of the ratios of critical stress for debonding c, and process zone radius, c for the SAN-24 and SAN-34 matrices at 80°C and 23°C=RT

Matrix c, 80 °C/c, RT c80 °C/cRT

SAN-24 0.66 1.32

SAN-34 0.62 1.25

20

List of Figure Captions

Fig. 1 AFM micrographs of SAN-24 containing (a) 0.10 vol.% Type-A, (b) 2.00 vol.% Type-A,

(c) 0.10 vol.% Type-B, (d) 0.50 vol.% Type-B, (e) 0.10 vol.% Type-C and (f) 1.25 vol.% Type-

C nano-particles. The arrows point to selected nano-particles while the circles indicate some

large sized agglomerates. Scale: Images are 2m x 2m.

Fig. 2 Typical DMTA traces showing the storage modulus and loss factor of SAN-24 containing 0.50 vol.% Type-B nano-rods tested at 1 Hz and 10 Hz.

Fig. 3 Normalised Young’s modulus versus particle volume fraction for the SAN-24

composites modified with Types A, B and C particles. Solid lines are best linear fits to the

normalised Young’s modulus and the horizontal dashed line represents zero enhancement of

stiffness. For well bonded particles at low volume fractions the slope is determined by Ep/Em,

the ratio of particle to matrix modulus, and expected to be 35. SAN-24-B follows this line at low

-values. Deviations would indicate interfacial failure (debonding) i.e. either particle-matrix or

particle-particle for agglomerates.

Fig. 4 Normalised fracture energy at 23 °C versus particle volume fraction for the SAN-24

composites modified with Types A, B and C particles. The horizontal dashed line represents

zero enhancement of fracture energy.

Fig. 5 Normalised fracture energy versus particle volume fraction for the SAN-24 and SAN-34 composites modified with Type-B nano-rods tested at 23 °C. The horizontal dashed line represents zero enhancement fracture energy.  

Fig. 6 Normalised fracture energy versus particle volume fraction for the (a) SAN-24 and (b) SAN-34 composites modified with Type-B nano-rods tested at 23 °C and 80 °C. Whilst fracture energy exhibits a maximum at quite low particle volume fractions at 23 °C a continuous increase is observed at a temperature closer to, but still below, the glass transition temperature of the matrix. The gain is highest for the more brittle SAN-24 matrix.

21

Fig. 7 SEM micrographs showing the fracture surfaces of: (a) unmodified SAN-24, (b) SAN-24 containing 0.05 vol.% Type-B nano-rods, (c) unmodified SAN-34 and (d) SAN-34 containing 0.10 vol.% Type B nano-rods tested at 23 °C (Crack propagation was from left to right). Arrows point to some voids and the circles indicate fibrillation of the matrix.

Fig. 8 SEM micrographs showing the fracture surfaces of: (a) unmodified SAN-24, (b) SAN-24 containing 0.10 vol.% Type-B nano-rods, (c) SAN-24 containing 0.50 vol.% Type-B nano-rods and (d) SAN-34 containing 1.00 vol.% Type-B nano-rods tested at 80 °C (Crack propagation was from left to right). Selected voids are indicated with arrows.

22

(b)(a)

(d)(c)

             A 0.1%                                                                            A  2%

  

  

                               B 0.1%                                                                    B 0.5%

continues onto next page

23

(c) (d)

500 nm500 nm

500 nm500 nm

(b)(a)

                       C   0.1%                                                                                 C     1.25%

  

Fig. 4 AFM micrographs of SAN-24 containing (a) 0.10 vol.% Type-A, (b) 2.00 vol.% Type-A,

(c) 0.10 vol.% Type-B, (d) 0.50 vol.% Type-B, (e) 0.10 vol.% Type-C and (f) 1.25 vol.% Type-

C nano-particles. The arrows point to selected nano-particles while the circles indicate some

large sized agglomerates.

24

(f)

500 nm500 nm

(e)

0 20 40 60 80 100 120 140 1600

1

2

3

4

5

0.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

Modulus @ 1 Hz

Modulus @ 10 Hz

tan(δ) @ 1 Hz

tan(δ) @ 10 Hz

Temperature, (°C)

Stor

age

mod

ulus

, E' (

GPa

)

Loss

fact

or, t

an(δ

)Fig. 5 Typical DMTA traces showing the storage modulus and loss factor of SAN-24 containing

0.50 vol.% Type-B nano-rods tested at 1 Hz and 10 Hz.

25

Fig. 6 Normalised Young’s modulus versus particle volume fraction for the SAN-24

composites modified with Types A, B and C particles. Solid lines are best linear fits

to the normalised Young’s modulus and the horizontal dashed line represents zero

enhancement of stiffness. For well bonded particles at low volume fractions the

slope is determined by Ep/Em, the ratio of particle to matrix modulus, and expected

to be 35. SAN-24-B follows this line at low -values. Deviations would indicate

interfacial failure (debonding) i.e. either particle-matrix or particle-particle for

agglomerates.

26

Fig. 4 Normalised fracture energy at 23 °C versus particle volume fraction for the SAN-24

composites modified with Types A, B and C particles. The horizontal dashed line represents

zero enhancement of fracture energy.

27

Fig. 5 Normalised fracture energy versus particle volume fraction for the SAN-24 and SAN-34

composites modified with Type-B nano-rods tested at 23 °C. The horizontal dashed line

represents zero enhancement fracture energy.

28

29

   (b)

Fig. 6 Normalised fracture energy versus particle volume fraction for the (a) SAN-24 and (b) SAN-34 composites modified with Type-B nano-rods tested at 23 °C and 80 °C. Whilst fracture energy exhibits a maximum at quite low particle volume fractions at 23 °C a continuous increase is observed at a temperature closer to, but still below, the glass transition temperature of the matrix. The gain is highest for the more brittle SAN-24 matrix.

30

 

 

Fig. 7 SEM micrographs showing the fracture surfaces of: (a) unmodified SAN-24, (b) SAN-24

containing 0.05 vol.% Type-B nano-rods, (c) unmodified SAN-34 and (d) SAN-34 containing

0.10 vol.% Type B nano-rods tested at 23 °C (Crack propagation was from left to right). Arrows

point to some voids and the circles indicate fibrillation of the matrix.

31

(b)(a)

(d)(c)

 

 

Fig. 8 SEM micrographs showing the fracture surfaces of: (a) unmodified SAN-24, (b) SAN-24

containing 0.10 vol.% Type-B nano-rods, (c) SAN-24 containing 0.50 vol.% Type-B nano-rods

and (d) SAN-34 containing 1.00 vol.% Type-B nano-rods tested at 80 °C (Crack propagation

was from left to right). Selected voids are indicated with arrows.

32

(b)(a)

(d)(c)