thesis bilotti 2009

Polymer / Sepiolite Clay

Nanocomposites

A THESIS SUBMITTED TO THE UNIVERSITY OF LONDON

FOR THE DEGREE OF DOCTOR OF PHILOSOPHY

February 2009

By

Emiliano Bilotti

School of Engineering and Materials Science

Queen Mary, University of London

Mile End Road, London E1 4NS

Declaration

I declare that the work presented in this thesis is performed entirely by myself during

the course of my PhD studies at Queen Mary, University of London and has not been

submitted for a degree at this or any other university.

Emiliano Bilotti

2

Acknowledgements

First of all, I would like to thank my supervisor Prof. Ton Peijs for giving me the

opportunity to carry out this PhD study in his group but also for his support, trust and

sense of humour. Secondly, I would like to thank Dr. Hartmut Fischer for his co-

supervision, for his hospitality at TNO but especially for his constant encouragement

and support throughout this study.

The vast majority of the experiments have been carried out at Queen Mary, Materials

Department, in a unique friendly atmosphere. I wish to thank all the group members

and support staff for many fruitful discussions and suggestions and more importantly

for their friendship. Thanks are especially given to Paola Ciselli, Chris Reynolds,

Hua Deng, Rui Zhang, Zhujuan Wang, Nattakan Soykeabkaew but also to Jianmin

Zhang, Saharman Gea, Antonio Scherillo, Manuela Russo, Giuseppe Viola, Marc

Simonet, Jia Ma, Luca Achilli, Dhanushka Hapuarachchi, Shuangwu Li, Wei Wang

and Chris Morgan. I would like to acknowledge Wenrui Zhang, Dun Lu, Michael

Dollinger, Franck Quero for direct cooperation and assistance with the experimental

work.

I wish to thank Monisha Phillips for her constant and valuable help, Zofia Luklinska

and Mick Wills for support on SEM and TEM and Vince Ford, Danny Neighbour

and Bill Godwin for manufacturing several experimental equipments.

I would like to thank Dr. Asa Barber for discussion, encouragement and support with

the SPM and Prof. Paul Smith and Prof. Bela Pukanszky for their critical comments

and suggestions. Prof. Philippe Dubois, Dr. Emmanuel Duquesne and Dr. Gaelle

Deshayes are kindly acknowledged for their hospitality at the University of Mons-

Hainaut and their guidance for the polymerisation of PA6/Sep hybrids.

3

Acknowledgements

Dr. Wim Brass and Dr. Kristina Kvashnina are thanked for the help and support with

WAXS experiments at the ESRF synchrotron facility.

Finally I would like to thank Claudio Bilotti for the experimental work on

electrospinning but even more for his constant enthusiasm and encouragement.

Inoltre, vorrei esprimere la più sincera gratitudine alla mia famiglia per il costante

supporto dimostrato durante questi anni. Un pensiero speciale va alla persona che più

mi è stata vicina in questa lunga avventura: Lydia. Grazie di cuore. Gli anni del PhD

sono stati la metafora della vita passata insieme fino ad oggi. Dalla nostra prima

conoscenza, con l’inizio del dottorato, all’amore durante il prosieguo fino al nostro

matrimonio celebrato pochi giorni fa.

4

Abstract

In the last two decades, polymer-clay nanocomposites have attracted great interests

because of the remarkable enhancements in mechanical and physical properties with

minute amount of nano-filler, promising to eliminate the typical compromise that

exists between properties and processability of composite materials. Despite the

expectations created by nano-clays in the academic and industrial communities, their

success has so far been limited. The reasons can be ascribed to the poor dispersion of

nano-clays in polymer matrices, to the often weak interfacial interaction with

polymers, and to the lack of control of nano-clay orientation.

In this thesis, all the aspects above will be tackled, studying the potential of sepiolite,

a nano-clay with a peculiar needle-like shape, in two thermoplastic polymers:

polypropylene (PP) and polyamide 6 (PA6). After an extensive literature survey, the

experimental part of the thesis starts with the characterisation of sepiolite (Chapter

5). The dimensions of the nano-filler are evaluated, as well as the specific surface

area, the thermal properties and the stiffness of individual nano-needles with novel

nano-mechanical tests. Chapter 6 deals with improving the dispersion and interaction

of nano-clay in PP, by employing three compatibilisers and by surface

functionalisation of the clay. In order to align the fibrous nano-filler, thin

PP/sepiolite tapes are prepared by solid-state drawing (Chapter 7). It is shown that in

these oriented nanocomposite tapes, small amounts of clay (<2.5wt%) are able to

improve the mechanical performances, due to a better reinforcement efficiency of the

nanofiller and to a higher orientation induced crystallinity, but also the thermal

resistance and barrier properties. Finally (Chapter 8-9), PA6/sepiolite

nanocomposites were prepared both by melt compounding and in situ

5

Abstract

polymerisation. Sepiolite is well dispersed in PA6 without any need of

compatibilisers or clay surface functionalisation. The simple and environmentally

friendly melt-compounding process turns out to be as efficient as the in situ

polymerisation route for what concerns the dispersion state of nano-filler, making it a

good candidate for industrial and commercial applications. In conclusion, sepiolite

shows promising credentials as nanofiller for thermoplastic polymers (i.e. PA6), and

in particular for oriented tapes, ultimately creating a 1D nanocomposite reinforced by

a 1D nanofiller.

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Table of Contents

List of Tables 12

List of Figures 14

1. Introduction 1.1 Nanoclay Composites - The present 25

1.2 Nanocomposites in Nature 27

1.3 Scope of the Thesis 30

1.4 References 30

PART 1: LITERATURE SURVEY

2. Clays and Clay Minerals 2.1 Introduction 33

2.2 Nomenclature of Clay Minerals 34

2.3 Sepiolite Clay 38

2.4 References 42

3. Polymer-Clay Nanocomposites 3.1 Introduction 43

3.2 Preparation of Polymer-Clay Nanocomposites 48

3.2.1 Solution Processing 48

3.2.2 In-Situ Polymerisation 48

3.2.3 Melt Processing 50

3.2.4 Template Synthesis 59

3.3 Properties of Polymer-Clay Nanocomposites 60

3.3.1 Crystallisation 60

3.3.2 Mechanical Properties 65

3.3.3 Barrier Properties 71

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Table of Contents

3.3.4 Fire Retardancy 74

3.4 References 76

4. Needle-like Clay Nanocomposites 4.1 Introduction - The importance of shape 86

4.2 Preparation of needle-like clay nanocomposites 88

4.2.1 In-situ polymerisation 88

4.2.2 Melt compounding 93

4.3 Properties of needle-like clay nanocomposites 94

4.3.1 Crystallisation 94


4.3.3 Rheology 99

4.4 References 102

PART 2: EXPERIMENTAL RESULTS AND DISCUSSION

5. Sepiolite Nanoclay in Polymer Composites 5.1 Introduction 106

5.2 Experimental 107

5.2.1 Materials 107

5.2.2 Characterisation Techniques 107

5.3 Results and Discussion 108

5.3.1 Morphological Analysis 108

5.3.2 B.E.T. Measurements 112

5.3.3 Thermal Properties 115

5.3.4 Mechanical Properties. Nano-Bending Tests 117

5.3.5 Micromechanical Models 120

5.4 Conclusions 130

5.5 References 130

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Table of Contents

6. Polypropylene / Sepiolite Nanocomposites 6.1 Introduction 134


6.2.1 Materials 135

6.2.2 Nanocomposites Preparation 135

6.2.2 Nanocomposites Characterisation 136



6.3.2 Crystal Structure and Crystallisation Behaviour 140

6.3.4 Rheological Behaviour 145

6.3.5 Thermal Behaviour 147



6.4 Conclusions 161

6.5 References 162

7. Oriented PP / Sepiolite Composite Tapes 7.1 Introduction 165


7.2.1 Materials 165

7.2.2 Composite Tapes Preparation 166

7.2.3 Composite Tapes Characterisation 167


7.3.1 Morphology of Tapes 168

7.3.2 PP and Sepiolite Orientation 170


7.3.4 Thermal Analysis 185

7.4 Conclusions 187

7.5 References 188

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Table of Contents

8. Polyamide 6 / Sepiolite Nanocomposites 8.1 Introduction 190


8.2.1 Materials 191

8.2.2 Nanocomposites Preparation 191




8.3.2 Non-Isothermal Crystallisation 195

8.3.3 WAXS – PA6 Crystal Structure 198

8.3.4 Thermo Gravimetric Analysis 203



8.4 Conclusions 213

8.5 References 214

9. In-Situ Polymerised Polyamide 6 / Sepiolite Nanocomposites 9.1 Introduction 217


9.2.1 Materials 217

9.2.2 Masterbatches Preparation - Polymerisation 218

9.2.3 Masterbatches Dilution - Melt Compounding 219



9.3.1 Gel Permeation Chromatography 221

9.3.2 TGA 223


9.3.4 Crystallisation Behaviour 228


9.3.6 Dynamic Mechanical Analysis 233

9.4 Conclusions 236

10

Table of Contents

9.5 References 237

10. Conclusions and Future Work

10.1 Summary 239

10.2 Future Work 242

10.3 References 245

List of Author’s Publications 247

11

List of Tables

1.1 Examples of commercial ventures in polymer/clay nanocomposites 26

2.1 Classification of clay minerals. From [1] 37

2.2 Proposed classification of phyllosilicate (as submitted by AIPEA

Nomenclature Committee to the International Mineralogical

Association). From [1] 38

2.3 Sepiolite physic-chemical properties on the basis of common

industrial applications 40

3.1 WAXS peak intensity and young’s modulus for nanocomposites

obtained by in-situ polymerisation, catalysed with different acids.

From [83] 65

3.2 Effect of PP-MA on the young’s modulus of PP-based

nanocomposites. From [54] 67

3.3 Effect of PP-MA on the yield stress of PP-based nanocomposites

[54] 69

3.4 Effect of PP-MA on the strain at break of PP-based


4.1 Mechanical properties of in-situ polymerised PE/palygorskite

nanocomposites [7] 97

4.2 Mechanical properties of rubbery epoxy/palygorskite


5.1 Volume of N2 gas adsorbed on sepiolite clay at different pressures 113

5.2 Weight losses of sepiolite clays 117

5.3 Results of the nano-bending tests 119

5.4 Shape factors ζ, for fibre-like and platelet-like reinforcement [16] 126

6.1 Compositions of PP/sepiolite nanocomposites 136

6.2 Relevant PP/smectite clay nanocomposites reported in the literature 154

8.1 Filler content of nanocomposites samples as obtained from TGA 204

8.2 Relevant PA6/clay nanocomposites reported in the literature 208

12

List of Tables

9.1 GPC results of molecular weight averages (Mw and Mn) and

polydispersity (Mw/Mn) of the commercial PA6, the two in-situ

masterbatches, and the nanocomposites obtained after dilution of

the masterbatches with commercial PA6. Samples are run in

duplicates 222

9.2 Monomer content in the in-situ polymerised masterbatches,

PA6/Sep and PA6/Sep-NH2, before and after extraction in hot water 225

9.3 Filler content of the different nanocomposites, calculated by the

residue of inorganic phase after TGA tests 226

9.4 Summary of the crystallinity and temperature of crystallisation of

the In-situ PA6 nanocomposites 229

13

List of Figures

1.1 Wilson's Double Core premium tennis ball, with technology Air D-

Fense, by InMat [3] 26

1.2 Hierarchical structure of bones: a) cortical and cancellous bone; b)

osteons with Haversian systems; c) lamellae; d) collagen fibre

assemblies of collagen fibrils; e) bone mineral crystals, collagen

molecules and non-collagenous proteins. From [4] 28

1.3 Structural hierarchy of the gecko adhesive system. (A)

Macrostructure: ventral view of a tokay gecko (G. gecko) climbing

vertical glass. (B) Mesostructure: ventral view of the foot, with

adhesive lamellae (scansors) visible as overlapping pads. Note the

clean appearance of the adhesive surface. (C) Microstructure:

proximal portion of a single lamella, with individual setae in an array

visible. (D and E) Nanostructure: single seta with branched structure at

upper right, terminating in hundreds of spatular tips [6]. 29

2.1 Electron microscope micrographs of different clays: a) kaolinite, b)

montmorillonite, c) halloysite and d) sepiolite. 34

2.2 Tetrahedral sheet 34

2.3 Octahedral sheet 35

2.4 Proposed crystallographic structures for: a) kaolinite, b) smectite and

c) chlorite. From [1] 36

2.5 Structure of sepiolite clays: a) SEM picture of natural sepiolite, b)

Schematic representation of a single sepiolite fibre, c) TEM picture

representing a single nanofibre cross-section and d) suggested

mineralogical structure. 39

3.1 Schematic representation of modification of platelet-like nano-clays

by cation-exchange reaction with surfactant molecules. From [10] 44

3.2 Interlayer structures of intercalated surfactant molecules into the nano-

clays galleries. From [11] 45

14

List of Figures

3.3 Alkyl chain aggregation model. As the chains length increase the

structure changes from liquid-like to liquid crystalline-like. From [11] 45

3.4 In situ SAXS-WAXS experiments during melting/recrystallisation.

The appearance of a small angle reflection after polymer melting

shows partial re-agglomeration of the clays. From [10] 46

3.5 Representation of the possible morphology arising from the interaction

of nanoclays and polymer: a) non-intercalated, conventional micro-

composite; b) intercalated nanocomposite and c) exfoliated

nanocomposite. From [4] 47

3.6 Representation of the swelling of nanoclays by ε-caprolactam and its

ring-opening polymerisation. From [5] 49

3.7 Schematic representation of melt-intercalation processing. From [5] 51

3.8 polymerVSh∆ , and as a function of the gallery height

, for a polymer and a layered silicate modified with

octadecylammonium group. Reproduced from [36]

chainVSh∆ VSh∆

0hh −

52

3.9 Variation of free energy per unit area, as a function of the

gallery height , for a polymer and a layered silicate modified

with octadecylammonium group. Curves are calculated for different

values of interaction energies ε

Vfh∆

0hh −

sp,sa between polymer and silicates,

supposing the interaction between polymer and tethered chains, εsa=0.

Curves I, IIa, IIb, III correspond to εsp,sa values of 0, -4, -8, -12 mJ/m2.

Reproduced from [36] 53

3.10 Representation of the “slurry” melt-compounding process. From [44] 55

3.11 Intercalation of maleic anhydride molecules into the clays galleries.

From [52] 57

3.12 Intercalation and successive exfoliation process of clays by melt-

compounding with PP-MA. From [5] 58

3.13 Polarised light micrographs of (a) neat PP (b) PP+4%clay

nanocomposites crystallised at 150°C after Yuan and Misra [73] 61

3.14 DSC traces of neat PP and PP+4%clay nanocomposites. Reproduced 61

15

List of Figures

by Yuan & Misra [73]

3.15 Hydrogen bonding in the α and γ crystalline forms of nylon 6 as seen

from end and side-view of each crystal. Closed and open circles

represent chain axes projecting out of and into the page, respectively.

From [76] 62

3.16 TEM pictures showing the crystalline morphology of a) PA6 and b)

nanocomposites with 2.5 wt.% of clays. Reproduced by Shen et al.

[81] 63

3.17 WAXS spectra of PA6 nanocomposites compared with pure PA6 (a) at

the surface and (b) in the bulk of injection moulded bars. PA2 and

PA3 refer to two different organo-modified MMT with the second

highly swollen. Reproduced by Varlot et al. [82] 64

3.18 Tensile modulus, measured at 120°C, of nanocomposites based on

MMT (● NCH) and Saponite (■ NCHP) clays at different loadings.

From [75] 66

3.19 Effect of clay content on tensile modulus of PA6-organomodified

MMT nanocomposites obtained by melt compounding. From [34] 66

3.20 Schematic formation of hydrogen bonds in PA6/MMT


3.21 Yield stress of PA6 nanocomposites for different MMT loadings.

Three molecular weights PA6 matrices are melt compounded with

organo-modified MMT clays. From [41] 68

3.22 Effect of clay content on Tensile Strength of co-intercalated PP/MMT


3.23 Elongation at break of three MW PA6 nanocomposites for different

MMT loadings, tested at crosshead spead of (a) 0.51 cm/min and (b)

5.1 cm/min. From [41] 71

3.24 Formation of a “tortuous path” in polymer-clay nanocomposites 72

3.25 Effect of exfoliation. Relative permeability plotted as a function of

aggregates width, for different sheet lengths. From [87] 73

3.26 Relative permeability for different loading of MMT clays. The best

16

List of Figures

fitting is for aspect ratio 192, much lower that 2000 expected for

completely exfoliated MMT platelets. From [87] 74

3.27 Heat release rate of PA6 and PA6/silicate nanocomposites (5 wt.%).

Reproduced from [89] 74

3.28 Residues of combustion of: (a) EVA with 5 phr organoclays; (b) EVA

with 5 phr MWCNTs; (c) EVA with 2.5 phr pure MWCNTs and 2.5

phr organoclays. From [92] 75

4.1 Surface area to volume ratio (A/V) as a function of the aspect ratio

(l/d) of cylindrical particles. Reproduced from [1] 86

4.2 Reinforcement effect of platelets and fibres in unidirectional

composites, for different aspect ratios, according to Halpin-Tsai and

Mori-Tanaka models. Reproduced from [3] 87

4.3 SEM micrographs of PA6/attapulgite nanocomposites obtained by in-

situ polymerisation. The pictures refer to filler concentration of: a) 2

wt.% and b) 5 wt.% (right). From [4] 88

4.4 TEM picture of boehmite stabilised in n-propanol. Reproduced from

[6] 89

4.5 TEM pictures of PA6/boehmite nanocomposites, obtained by in-situ

polymerisation, referring to (a) 7.5 wt.% and (b) 9 wt.% of filler. From

[6] 90

4.6 TEM pictures of two different concentrations of Ti-modified

boehmite/PA6 nanocomposites: 7 wt.% (left) and 15 wt.% (right).

From [5] 90

4.7 TEM picture of sepiolite in epoxy matrix. From [8] 91

4.8 TEM pictures of elastomer/sepiolite nanocomposites containing 5phr

of clays. From [10] 92

4.9 TEM picture of PP/org-attapulgite nanocomposites with 5 wt.% filler

content [18] 94

4.10 Avrami’s plot for (a) neat PP and PP/ATP nanocomposites: (b)

PP/ATP 1 wt.%, (c) PP/ATP 3 wt.% and (c) PP/ATP 5 wt.%, at three

different crystallisation temperatures [21] 96

17

List of Figures

4.11 Young’s modulus and yield stress for PP/Org-ATP nanocomposites in

function of the filler content [19] 97

4.12 Small amplitude strain sweep (A) and frequency sweep (B-D) at 260

°C for neat PA6 (a) and PA6 nanocomposites with Attapulgite content

of: (b) 2 wt.%, (c) 3 wt.%, (d) 4 wt.% and (e) 5 wt.%. From [4] 100

4.13 Schematic description of the polymer/needle-like clay percolating

structure. From [22] 101

4.14 2D sketch of the percolation lattice model at Φ<Φc (A), Φ=Φc (B) and

Φ>Φc. Black occupied lattices represent the sticks percolated, while

the grey occupied lattices represent the sticks unpercolated. The

symbol ‘X’ indicates the lattices occupied by grafted polymer chains.

From [4] 101

5.1 SEM micrograph of sepiolite clays dispersed on a porous substrate 109

5.2 TEM micrograph of sepiolite on carbon-coated copper grids 109

5.3 TEM micrographs of sepiolite clays on carbon coated TEM grids.

Distribution of: a) lengths and b) diameters. Black arrows underline

single fibres measurements 110

5.4 Distributions of: a) sepiolite lengths and b) sepiolite diameters 111

5.5 B.E.T. plot of )/1()/( 0PPVPP ado − versus relative pressure .

C and can be calculated from the linear fit of the data point, since

the slope is and intercept is

)/( oPP

admonV

admonCVC /)1( − ad

monCV/1 114

5.6 Dehydration of sepiolite clays under temperature scan. The dashed

lines represent the temperature window at which clays are typically

subjected during composites preparation (extrusion and compression

moulding) 116

5.7 Schematic representation of SPM bending test of sepiolite nano-fibres

suspended on porous substrate 118

5.8 SPM tests: a) image of a sepiolite nanoclay laying on the substrate and

b) typical force-displacement curve 118

5.9 Schematic illustration of the concept of the shear lag model: a)

18

List of Figures

unstressed and stressed system and b) variation of the shear stress and

strain in the matrix in function of the radial position. From [25] 120

5.10 Predicted variations in a) fibre tensile stress and c) interfacial shear

stress along the length of a glass fibre (schematically represented in

b)), in polyester/30 % glass fibre composite, subject to an axial tensile

strain of 10-3, for two fibre aspect ratios. Redrawn from [25] 122

5.11 Critical aspect ratios in function of the sepiolite vol.% for two polymer

composites (solid lines). The dashed horizontal lines represent the

average aspect ratio (middle) and the lowest and highest values of the

aspect ratio distribution (bottom and top), measured from TEM

micrographs 124

5.12 Principal directions of composites relative to the ones of oriented

fillers 126

5.13 Reinforcement of 5 vol.% of fibre-like and platelet-like filler,

unidirectionally oriented (1D) in two polymer composites. The dashed

vertical line shows the average sepiolite aspect ratio 128

5.14 Reinforcement of 5 vol.% of 3D randomly oriented fibre-like and

platelet-like fillers in PP matrix. The dashed line shows the average

sepiolite aspect ratio 129

6.1 SEM micrographs of: a) PP+2.5%Sep; b) PP+PP-g-MA+2.5%Sep; c)

PP+PP-acid+2.5%Sep; d) PP+1%Sep-sil and e) PP+5%Sep-sil. White

circles underline sepiolite clusters. A significant improvement in the

dispersion of sepiolite in PP matrix is evident with the use of PP-PEO

and Sep-sil, where no agglomerates of nanoclay are found in

nanocomposites at 2.5 wt.% filler load 139

6.2 TEM picture of: a) sepiolite dispersion on TEM grids b) sepiolite in

PP matrix after compounding. A reduction in fibre length is evident in

the processed nanocomposites as a consequence of melt blending in

mini-extruder 140

6.3 X-ray diffraction spectra of: a) PP+Sep; b) PP+PP-g-MA+Sep; c)

PP+PP-acid+Sep; d) PP+Sep-sil nanocomposites at different

19

List of Figures

concentrations of filler, compared with virgin PP and pure sepiolite 141

6.4 DSC traces corresponding to the non-isothermal crystallisation of a)

PP+PP-g-MA+Sep, b) PP+PP-acid+Sep at different filler content. The

exothermic peaks shift towards higher temperature as a result of the

filler nucleating effect 142

6.5 Onset Temperatures of starting crystallisation in function of the filler

concentration for: ■ PP+Sep; ○ PP+PP-g-MA+Sep; PP+PP-

acid+Sep; PP+Sep-sil. A larger and continuous increase in the

crystallisation temperature is observed for PP+Sep and PP+PP-g-

MA+Sep while a limiting concentration of crystallisation nuclei is

reached at 1 wt.% of filler for PP+PP-acid+Sep and PP+Sep-sil 143

6.6 Frequency sweep test on polypropylene at 200°C 145

6.7 Complex viscosity of PP/Sep-sil nanocomposites in function of the

clay loading 146

6.8 Storage modulus of PP/Sep-sil nanocomposites in function of the clay

loading 147

6.9 TGA of PP and PP+Sep-sil nanocomposites with different amounts of

nanoclays, in N2 148

6.10 TGA of PP and PP+Sep-sil nanocomposites with different amounts of

nanoclays, in air 149

6.11 Stress-strain curves of different nanocomposites with 1 wt.% of

sepiolite 150

6.12 Stress-strain curves of different nanocomposites with 5 wt.% of

sepiolite 151

6.13 Young’s modulus of PP nanocomposites at different filler loadings 151

6.14 Yield stress of PP nanocomposites at different filler loadings 152

6.15 Strain at break of PP nanocomposites at different filler loadings. While

nanocomposites with pristine clay and with PP-g-MA undergo a clear

embrittlement, the use of PP-acid and Sep-sil preserves ductility even

at filler concentrations above 5 wt.% 153

6.16 Young’s modulus of PP/clay nanocomposites in function of the filler

20

List of Figures

wt.% 155

6.17 Relative Young’s modulus versus filler vol.%. The lines are prediction

from the Halpin-Tsai equations for PP/Sep nanocomposites (fibre-like

filler; solid line) and PP/smectite clay nanocomposites (plate-like

filler; dotted line), using true filler aspect ratios (s) as a fitting

parameter. In the graph the prediction for fibre-like and plate-like filler

for s ∞→ , corresponding to the rule of mixtures, are also included.

The abbreviations ‘1D’ and ‘2D’ stand for uniaxially oriented and 2D

in-plane randomly distributed filler, respectively 156

6.18 Tensile yield stress of PP/clay nanocomposites in function of the filler

vol.% 159

6.19 The natural logarithm of relative tensile stress of PP/clay

nanocomposites in function of the filler volume percent. From the

linear fit of the experimental data, the parameter B can be extracted 160

7.1 Schematic illustration of nanocomposite tape preparation. A

rectangular specimen is cut from a 100 µm thick compression moulded

film and drawn in the solid state to a tape of the desired draw ratio 166

7.2 SEM micrographs, at different magnifications, of the lateral surface

of: a-b) PP tapes, c-d) PP+1%Sep tapes and e-f) PP+10%Sep, after

ductile failure. A white circle indicates sepiolite agglomerations 169

7.3 Scheme of WAXS measurements in the through direction 170

7.4 2D WAXS patterns of: a-b) PP, c-d) PP+5%Sep and e-f) PP+10%Sep

tapes, at λ=1 (left column) and λ=20 (right column), respectively.

Black arrows indicate the main reflection planes of PP and of sepiolite

clays 172

7.5 X-ray intensity (integrated along the 2θ axis) versus the Azimuth angle

for a) PP sample and b) PP+5%Sep tapes at λ=1, 7 and 20. Solid line

represents Gaussian fitting of the data points 174

7.6 Hermans’ orientation factor as a function of draw ratio λ, relative to

PP tapes with different concentrations of sepiolite 176

7.7 Hermans’ orientation factor for PP+5%Sep composite tape, as a

21

List of Figures

function of draw ratio λ 177

7.8 Stress-strain curves of PP tapes of different draw ratios 178

7.9 Young’s modulus of nanocomposites tapes at different draw ratios.

The dotted lines are Halpin-Tsai predictions of PP tapes filled with 2.5

wt.% sepiolite, completely aligned in the direction of the tape, at three

aspect ratios: s=12, which was found to fit the isotropic samples (Fig.

6.15), s=27, which is the average aspect ratio of sepiolite nanofibres

(5.3.1), and s , which corresponds to the upper bound of the rule

of mixtures

∞→

179

7.10 Young’s modulus of nanocomposite tapes of λ=9, 16 and 20, as a

function of sepiolite filler content 181

7.11 Ultimate tensile strength of nanocomposites tapes in function of λ 181

7.12 Ultimate tensile stress of nanocomposites tapes in function filler wt.%. 182

7.13 Strain at break of nanocomposites tapes 182

7.14 Degree of crystallinity of different nanocomposite tapes in function of

λ 186

7.15 Young’s modulus of nanocomposites tapes in function of the degree of

polymer crystallinity 187

8.1 SEM micrographs of: a)-b) PP+1%Sep, c)-d) PP+2.5%Sep, e)-f)

PP+5%Sep, at magnification of 5000 and 10000 times respectively. A

good dispersion of sepiolite nanoclays is evident even at relative high

filler content 193

8.2 TEM micrographs of PA6/sepiolite nanocomposites with 5 wt.% of

filler at different magnifications 194

8.3 DSC crystallisation peaks of PA6/sepiolite nanocomposites at

different filler concentrations 195

8.4 DSC melting peaks of PA6/sepiolite nanocomposites at different filler

concentrations 196

8.5 Amount of crystalline phase for PA6/sepiolite nanocomposites. The

heat of fusion for the completely crystalline PA6 is taken as 240 J/g

[3]. Virgin PA6 is presented, for comparison, as open circle, while full

22

List of Figures

squares refers to processed samples 197

8.6 Scheme of WAXS measurements in the through direction 198

8.7 WAXS through view: a) PA6, b) PA6+1%Sep, c) PA6+5%Sep. White

arrows show the principal diffraction rings with the corresponding

crystal planes 199

8.8 Scheme of WAXS measurements in the edge direction 200

8.9 WAXS, edge view: a) PA6, b) PA6+1%Sep and c) PA6+5%Sep.

White arrows shows orientation of sepiolite, as can be seen from the

diffraction at 2θ ~ 7.2 °. 201

8.10 Schematic 3D image of a nanocomposite tensile test specimen where

the nanofiller is aligned in-plane. Sepiolite nano-fibres are represented

in red and are not in scale with the specimen dimensions. The real

length of sepiolite clay is about 4000 times smaller than in the picture 202

8.11 TGA of PA6/Sep nanocomposites, in inert atmosphere (N2) 203

8.12 Stress-strain curves of PA6/Sep nanocomposites at different filler

loadings 204

8.13 Young’s modulus of PA6/sepiolite nanocomposites in function of the

filler loading 205

8.14 Ultimate tensile stress of PA6/sepiolite nanocomposites in function of

the filler loading 205

8.15 Strain at break for PA6/sepiolite nanocomposites in function of the

filler loading 207

8.16 Relative Young’s modulus versus filler vol.%. The lines are prediction

from the Halpin-Tsai equations for 2D randomly oriented PA6/Sep

nanocomposites (fibre-like filler; solid line) and uniaxially (1D)

oriented PA6/MMT nanocomposites (plate-like filler; dotted line),

using true filler aspect ratios (s) as a fitting parameter. The condition

corresponds to the upper bound predictions of the rule of

mixtures

∞→s

209

8.17 Tensile stress of PA6/clay nanocomposites in function of the filler

vol.% 211

23

List of Figures

8.18 The natural logarithm of relative tensile stress of PA6/Sep and

PA6/MMT nanocomposites in function of the filler vol.%. From the

linear fit of the experimental data, the parameter B can be extracted 212

9.1 TGA of In-situ PA6/Sep-NH2 masterbatch as produced and after three

successive extractions in formic acid (from bottom to top) 223

9.2 TGA in inert atmosphere (N2) of In-Situ PA6/20%Sep-NH2

masterbatch, before (solid line) and after (broken line) purification in

hot distilled water 224

9.3 SEM micrographs of: a)-b) In-situ PA6+5%Sep and c)-d) In-situ

PA6+5%Sep-NH2, at magnifications respectively of 10000 (left

column) and 50000 times (right column) 227

9.4 TEM micrographs of In-situ PA6+5%Sep-NH2 228

9.5 Stress-strain curves of: a) In-situ PA6/Sep and b) In-situ PA6/Sep-NH2

nanocomposites, at different nominal filler loadings 230

9.6 Elastic Moduli for In-Situ PA6/Sep and PA6/Sep-NH2

nanocomposites at different filler loadings 231

9.7 Ultimate tensile stress for In-situ PA6/Sep and PA6/Sep-NH2

nanocomposites at different filler loadings 231

9.8 Strain at break for In-Situ PA6/Sep and PA6/Sep-NH2 nanocomposites

at different filler loadings 232

9.9 Toughness of In-situ PA6 / Sep and PA6 / Sep-NH2 nanocomposites at

different filler loadings, calculated from the integration of the

engineering tensile tests curves over the strain 233

9.10 DMA: storage modulus and tan δ curves of PA6 234

9.11 Tg of In-situ PA6/Sep and In-situ PA6/Sep-NH2 nanocomposites vs.

filler loadings 235

10.1 TEM micrographs of: a)-c) HPC/PEO electrospun fibres filled with

sepiolite needle-like clay, d) aligned electrospun fibres and e)

electrospun fibres twisted into a yarn 244

24

1 Introduction

1.1 Nanoclay Composites - The present

In recent years polymer/clay nanocomposites have attracted great interest in

academia and not exclusively. Important enterprises are currently racing to

commercialise nanoclays thermoplastic composites in particular focusing on

automotive parts and packaging. The attractiveness of this new class of material lays

on the large improvements in the mechanical and thermal properties, as well as gas

barrier and flame resistance, provided by only small amounts of nanometre-size clays

homogeneously dispersed in a polymeric matrix. The first commercial

nanocomposites product was an engine cover-belt developed by Japan’s Ube

Industries after being licensed the patented in situ polymerised PA6/Clay

nanocomposites (NCH) technology by Toyota Research Centre, following the

pioneering work done about one decade ago [1-2]. Compared with unfilled nylon 6,

Ube's NCH was claimed to have 68% higher tensile modulus and 126% higher

flexural modulus, along with a reduction in oxygen permeability of 50%. More

recently, Bayer AG (Germany) announced the development of nylon 6

nanocomposites for transparent barrier film packaging. The enhancement in barrier

properties is the key of another application of nanocomposite. Wilson's Double Core

premium tennis ball (Fig.1.1), officially selected for the 2002 Davis Cup, is said to

retain its original air pressure and extend the life time of tennis balls. This is due to

25

CHAPTER 1 – Introduction

the presence of a flexible and very thin (10-50µm) nanocomposite barrier coating

(called Air D-Fense, by InMat) that covers the inner core and that inhibits air

permeation through the walls of the ball by a factor of two.

Figure 1.1. Wilson's Double Core premium tennis ball, with technology Air D-Fense,

by InMat [3].

The nanocomposite coating consists of well exfoliated vermiculite nano-clays,

oriented along the radial direction of the ball and embedded in a matrix of butyl

rubber. These platelets act as multi-layer barriers to the diffusion of air and accounts

for the better performances of the tennis balls.

For what concerns the investigations of other properties it is worth to mention the

announcement of a consortium of government and industry scientists, formed by The

National Institute of Standards and Technology (NIST) in Gaithersburg, Md. (USA),

to explore nanocomposites' potential for reducing the flammability of thermoplastics.

In Table 1.1 a series of commercial ventures in polymer-clay nanocomposites are

presented.

26


Table 1.1. Examples of commercial ventures in polymer/clay nanocomposites. UBE Nylon6 Toyota timing belt cover, engine manifold cover

Nylon6 Film for packaging

Nylon6/66, 12 Fuel system components

BAYER PLASTICS Nylon6 Film for meat packaging

Nylon6 coating for paper board juice container

PC/ABS Flame retardant computer and monitor housings

FORSTER CORP. Nylon 12 nanocomposites used in catheter tubing

GM Polyolefin TPO for step on Astro vans to replace talc filled materials

UNITIKA Nylon6 automotive parts (Mitsubishi engine cover)

EVOH, Polylactic acid nanocomposites (various automotive uses)

WILSON SPORTING Tennis balls (butyl rubber/nanoclay coating from InMat)

HONEYWELL Nylon6 for food packaging

U.S. ARMY MRE food tray (EVOH)

KABLEWERK EUPEN EVA flame retardant cable coating

TNO Polyacrylate binding system for ceramic moulds

MITSUBISHI Polypropylene nanocomposites for automotive parts

TRITON SYSTEM Polyurethane bladder for athletic shoes

Polyolefin packaging film for food and pharmaceutical packaging

NANOCOR MXD-6 Nylon for barrier food packaging

1.2 Nanocomposites in Nature

If we are just appreciating the properties that nano-fillers as nanoclays can guarantee,

Mother Nature already commonly uses nanocomposites and nano-scale design,

obtaining materials with extraordinary properties. The best way to understand what

might be the potentiality of nanocomposites is simply to observe Nature.

A very common example is the structure of bones. A bone has a complex

arrangement of materials and structures at different length-scales, which work in

synergy to perform diverse mechanical, biological and chemical functions, such as:

structural support, protection and storage of healing cells, and mineral ion

homeostasis. The structure of bone can be described as a hierarchical organisation

27


[4]. The different levels of this structure are: (1) the macrostructure: cancellous and

cortical bone; (2) the microstructure (from 10 to 500 µm): Haversian systems,

osteons, single trabeculae; (3) the sub-microstructure (1–10 µm): lamellae; (4) the

nanostructure (from a few hundred nanometres to 1 µm): fibrillar collagen and

embedded mineral; and (5) the sub nanostructure (below a few hundred nanometres):

molecular structure of constituent elements, such as mineral, collagen, and non-

collagenous organic proteins (Fig.1.2). The structure is made more complex by the

3D arrangement and orientations of the different components.

Figure 1.2. Hierarchical structure of bones: a) cortical and cancellous bone; b)

osteons with Haversian systems; c) lamellae; d) collagen fibre assemblies of

collagen fibrils; e) bone mineral crystals, collagen molecules and non-collagenous

proteins. From [4].

Another example of a hierarchically designed material, which is composed by

fundamental units on the nanometre scale, and that can achieve amazing

performances is the Gecko toe pads. Over 2000 years ago, Aristotle [5-6] commented

on the capacity of geckos to “run up and down a tree in any way, even with the head

downwards”. Although some mechanisms on how the Gecko can climb vertical flat

surfaces so easily are still not perfectly understood, we know that the reason lays on

28


the hierarchical structure of its toe pads that effectively functions as a smart adhesive

[6]. The adhesive lamellae on the toe pad are composed by micro-scale arrays of

setae (Fig.1.3.B-C) and each seta presents hundreds of nano-scale spatular tips at its

end (Fig.1.3.D-E). The result of this peculiar design is that the gecko’s toe: a)

attaches strongly with minimal preload, b) detaches quickly and easily [7], c) sticks

to nearly every material, d) does not “stay dirty” [8] or e) self-adhere, and 7) is non-

sticky by default.

Figure 1.3. Structural hierarchy of the gecko adhesive system. (A) Macrostructure:

ventral view of a tokay gecko (G. gecko) climbing vertical glass. (B) Mesostructure:

ventral view of the foot, with adhesive lamellae (scansors) visible as overlapping

pads. Note the clean appearance of the adhesive surface. (C) Microstructure:

proximal portion of a single lamella, with individual setae in an array visible. (D and

E) Nanostructure: single seta with branched structure at upper right, terminating in

hundreds of spatular tips [6].

29


1.3 Scope of the Thesis

Most of the literature in the field of Polymer/Clay nanocomposites is focused on

platelet-like clays, commonly smectite clays such as Montmorillonite. Few works

have instead been dedicated to fibre-like clays particles. Because of the peculiar

shape, these nano-fillers are believed to be good candidates for the preparation of

nanocomposites materials. In fact the dispersion of needle-like clays, compared to

platelet-like clays, is favoured by the relatively small contact surface area.

Furthermore the reinforcement efficiency of fibres is higher than platelet for uniaxial

composites.

The focus of this research is the investigation of sepiolite, a natural needle-like clay,

as a reinforcement for thermoplastic polymers. Two polymer matrices will be taken

in consideration: polypropylene (PP) and polyamide 6 (PA6).

Particular emphasis will be given to the nanocomposites preparation, improvements

in the inorganic filler dispersion, interphase compatibility and, not least,

nanocomposites characterisation and structure-properties relationship. Whenever

possible, the polymer/sepiolite nanocomposites prepared will be compared and

benchmarked with the more widely studied smectite clays nanocomposites from the

scientific literature, throughout this thesis.

1.4 References

1. Okada A, Kawasumi M, Usuki A, Kojima Y, Kurauchi T, Kamigaito O. Synthesis

and properties of nylon-6/clay hybrids. In: Schaefer DW, Mark JE, editors. Polymer

based molecular composites. MRS Symposium Proceedings, Pittsburgh, vol. 171;

1990. p. 45–50.

2. Usuki, A.; Kawasumi, M.; Kojima, Y.; Okada, A.; Kurauchi, T.; Kamigaito, O. J

Mater Res 1993, 8, 1174.

3. http://www.wilsonsports.com.au/tennis/doublecore.html

30


4. J. Rho, L. Kuhn-Spearing, P. Zioupos. Mechanical properties and the hierarchical

structure of bone. Medical Engineering & Physics. 20, 2 (1998) 92-102.

5. Aristotle (350 B.C.E., 1918) Historia animalium translated by Thompson, D'A-W.

Clarendon .Press, Oxford.

6. K. Autumn. Properties, Principles, and Parameters of the Gecko Adhesive System.

From: Biological Adhesives (ed. by AM. Smith and JA. Callow) 63 Springer-Verlag

Be& Heidelberg 2006.

7. K. Autumn, A. Peattie. Mechanisms of adhesion in geckos. Int. Comp Bio 42

(2002) 1081-1090.

8. W. Hansen, K. Autumn. Evidence for self-cleaning in gecko setae. PNAS, (2005)

102385-389.

31

PART 1: Literature Survey

32

2 Clays and Clay Minerals

2.1 Introduction

The term clay has often been defined operationally. According to Grim [1, 2] “the

term clay implies a natural, earthy, fine-grained material which develops plasticity

when mixed with a limited amount of water. By plasticity is meant the property of

the moistened material to be deformed under the application of pressure, with the

deformed shape being retained when the deforming pressure is removed”.

Clays are the main constituents of the fine-grained sedimentary rocks as mudstones

and shales in marine sediments and in soils and are the results of weathering and

secondary sedimentary processes with only a few examples of clays forming in

primary igneous or metamorphic environments.

In geology, the term clay includes particles <2 µm in size; the morphology of the

clay-mineral components being a distinctive property of a particular clay. For

instance, kaolinite usually shows hexagonal flake-shaped unites with a ratio of areal

diameter to thickness (aspect ratio) of 2-25:1, while most of smectite mineral

particles have an irregular flake shape but with a much higher aspect ratio, 100-

300:1. Halloysite minerals show an elongated tubular shape, while the family of

attapulgite/sepiolite/palygorskite are characterised by a peculiar elongated lath or

fibre-shape.

33

CHAPTER 2 – Clays and Clay Minerals

a) b)

c)

d)

Figure 2.1. Electron microscope micrographs of different clays: a) kaolinite, b)

montmorillonite, c) halloysite and d) sepiolite.

2.2 Nomenclature of Clay Minerals

Clay minerals belong to the family of phyllosilicates (or layered silicate). The

fundamental building units of phyllosilicates (and then of clay minerals) are

tetrahedral and octahedral sheets. Tetrahedral sheets are composed of individual

tetrahedrons, in which a silicon atom (but also Al3+, Fe3+, etc) is equidistant from

four oxygens, or hydroxyls if needed to balance the structure. They are arranged in a

hexagonal pattern with the basal oxygens linked and the apical oxygens pointing

up/down and taking part in the adjacent octahedral sheet.

Figure 2.2. Tetrahedral sheet

34


Octahedral sheets are composed of individual octahedrons that share edges

composed of oxygen and hydroxyl anion groups coordinated by cations like Al, Mg,

Fe3+ and Fe2+, etc.

Figure 2.3. Octahedral sheet

According to the valence of the cation we can distinguish di-octahedral or tri-

octahedral sheet, which structure resemble respectively the minerals Gibbsite

Al(OH)3 and Brucite Mg(OH)2. When we have a trivalent cation (i.e. Al3+), in order

to maintain electric neutrality, the cation to oxygen ratio is 1:3. This leaves every

third site empty, meaning only 2 out of 3 sites are occupied. This arrangement is

called di-octahedral or Gibbsite-like sheet. Instead, when we have a divalent cation

(i.e. Mg2+) occupying the edge sharing hexagonal sheet the cation to oxygen ratio is

1:2 and every lattice site is filled. This arrangement is called tri-octahedral or

Brucite-like sheet.

The main criterion of classification of phyllosilicates is determined by the way the

different tetrahedral, di- and tri-octahedral sheets are packed together. The structure

of kaolinite, for instance, is composed by one silica tetrahedral sheet and one alumina

octahedral (1:1) sheet combined to form a layer unit, in which the apical oxygens of

the tetrahedral sheet are also part of the octahedral sheet (Fig. 2.4.a). Smectite is

composed of units made up of two silica tetrahedral sheets with a central alumina

octahedral sheet (2:1) (Fig. 2.4.b). The structure of chlorite, instead, can be imagined

as consisting of alternating smectite-like layers and a brucite-like tri-octahedral sheet

(Fig. 2.4.c).

In clay minerals, except for kaolinite, a certain number of cations are replaced by

ions of lower valence. So Si3+ in the tetrahedral sheet can be replaced by Al2+, while

Al2+ in the octahedral sheet may be replaced by Li+, Mg2+, Fe2+, Fe3+, Zn2+, etc.

35


These isomorphous substitutions, along with the presence of vacancies, account for a

negative charged surface of the clays layers. The net negative charge, often denoted

as CEC (Cation Exchange Capacity) and expressed as mequiv/100 g, are

counterbalanced by alkali and alkaline earth cations situated inside the galleries

(defined as the space between two layer units).

a)

b)

c)

Figure 2.4. Proposed crystallographic structures for: a) kaolinite, b) smectite and c)

chlorite. From [1].

In the case of tetrahedrally substituted layered silicates, the negative charge is located

on the surface of silicate layers, and hence, it is more ready for interaction (for

instance with polymers) compared with octahedrally substituted material. The

interlayer charge can create a bonding between different layers also relatively strong

36


that are closely packed and difficult to exfoliate. In montmorillonite (MMT), on the

other hand, the ions are exchangeable, the distance between the layers can thereby

increase and the material can swell (denoted as swelling clay minerals). Two

different classifications of clay minerals are presented in Table 2.1-2.2.

Table 2.1. Classification of clay minerals. From [1].

I. Amorphous Allophane group

II. Crystalline A. Two-layer type (sheet structures composed of units of one layer of silica

tetrahedrons and one layer of alumina octahedrons. 1. Equidimensional

Kaolinite group Kaolinite, Nacrite, etc.

2. Elongate Halloysite group

B. Three-layer types (sheet structures composed of two layers of silica tetrahedrons and one central di-octahedral or tri-octahedral layer)

1. Expanding lattice a. Equidimensional

Montmorillonite group Montmorillonite, sauconite, etc. Vermiculite

b. Elongate Montmorillonite group Nontronite, Saponite, hectorite

2. Nonexpanding lattice Illite group

C. Regular mixed-layer types (ordered stacking of alternate layers of different types)

Chlorite D. Chain-structure type (hornblende-like chains of silica tetrahedrons linked

together by octahedral groups of oxygens and hydroxyls containing Al and Mg atoms)

Attapulgite Sepiolite Palygorskite

37


Table 2.2. Proposed classification of phyllosilicate (as submitted by AIPEA

Nomenclature Committee to the International Mineralogical Association). From [1].

Type Group (x=layer charge)

Subgroup Species

Pyrophyllites Pyrophyllite Pyrophyllite-talc x~0 Talcs Talc

Dioctahedral smectites or montmorillonite

Montmorillonite, beidellite, nontronite

Smectite or montmorillonite-saponite x~0.5-1

Trioctahedral smectites or saponites

Saponite, hectorite, sauconite

Dioctahedral vermiculite Dioctahedral vermiculite Vermiculite x~1-1.5 Trioctahedral vermiculite Trioctahedral vermiculite

Dioctahedral micas Muscovite, paragonite Mica x~2 Trioctahedral micas Biotite, phlogopite

Dioctahedral brittle micas Margarite

2:1

Brittle mica x~4 Trioctahedral brittle micas Seybertite,

xanthophyllite, brandisite Dioctahedral chlorite 2:1:1 Chlorite

x variable Trioctahedral chlorite Pennine, clinochlore, prochlorite

Kaolonites Kaolinite, halloysite 1:1 Kaolinite-serpentine x~0 Serpentines Chrysotile, lizardite,

antigorite

2.3 Sepiolite Clay

Sepiolite is a fibrous hydrated magnesium silicate, a typical formula for which is

Mg4Si6O15(OH)2·6H2O. The name comes from a perceived resemblance of the

material to the porous bones of the cuttlefish or sepia. Sepiolite is included in the

phyllosilicate group because it contains a continuous two-dimensional tetrahedral

sheet of composition Si O [1, 3]. It differs, however, from the other layered silicates

because of the lack of a continuous octahedral sheet (Fig. 2.5.d). 2 5

It can be imagined

as formed of blocks structurally similar to layered clay minerals (i.e. MMT),

composed of two tetrahedral silica sheets and a central octahedral sheet containing

Mg, but continuous only in one direction (c-axis). More blocks are linked together

along their longitudinal edges by Si-O-Si bonds and this creates channels along the c-

38


axis (Fig. 2.5.b-d). Moreover, because of the covalent link between different blocks,

sepiolite has been described as a non-swellable clay. Due to the discontinuity of the

external silica sheets, a significant number of silanol groups (SiOH) are situated at

the edges of this mineral.

a)

b)

c)

d)

Figure 2.5. Structure of sepiolite clays: a) SEM picture of natural sepiolite, b)

schematic representation of a single sepiolite fibre, c) TEM picture representing a

single nanofibre cross-section and d) suggested mineralogical structure.

Sepiolite is composed by elemental particles with needle-like of fibre-like shape. The

dimensions of the a single sepiolite fibre vary between 0.2-4µm in length, 10-30nm

in width and 5-10nm in thickness, with open channels of dimensions 3.6 Å x 10.6 Å

running along the axis of the particle (Fig. 2.5.a-b). These particles are arranged

forming loosely packed and porous aggregates with an extensive capillary network

which explains the high porosity. Sepiolite has the highest surface area of all the clay

minerals, about 300 m2/g and a high sorption capacity. There are three sorption sites:

(a) oxygen ions on the tetrahedral sheet, (b) a small amount of cation-exchange sites

(0.1-0.6 mequiv/100 g) and (c) the already mentioned SiOH groups. Adsorption is

also influenced by the size, shape and polarity of the molecules involved. Neither

large molecules nor those of low polarity can penetrate the channels though they can

39


be adsorbed on the external surface, which accounts for about 50-60 % of the total

surface area [4, 5]. The SiOH groups act as neutral sorption sites for suitable for

organic species. Apart from the outstanding sorptive capacity, sepiolite is also known

for its colloidal properties. When dispersed in a liquid, it forms a structure of

randomly intermeshed elongated particles, which is maintained by secondary bonds.

This structure is stable even in systems with high salt concentrations, condition that

produces the flocculation of other clay’s suspensions, as bentonite. Sepiolite provides

a pseudoplastic and thixotropic behaviour which make it a valuable material in

multiple applications to improve processability, application or handling of the final

product. Common industrial applications of sepiolite are listed in Table 2.3.

Table 2.3. Sepiolite physic-chemical properties on the basis of common industrial

applications.

Application Characteristic

Cat and pet litters Light weight, high liquid absorption, odour control.

Industrial absorbents High liquid absorption, mechanical strength in wet

conditions, non-flammability, chemical inertness.

Carrier for chemicals Absorption of active chemicals and easiness and

effectiveness in delivering them.

Bitumens Control of rheological properties in heat application

systems, improving fire resistance.

Rheological additives Stability, pseudo-plasticity and thixotropy in paints,

adhesives, mastics and sealants.

Health and Safety

The use of nanoparticles has recently raised several health issues. It is then

fundamental to understand the risks in handling sepiolite and using it as a nanofiller

for polymeric matrices. The health and safety assessments of the sepiolite used

(River Tajo basin, Madrid, Spain) supplied by Tolsa (Spain), including

40


epidemiological, in vitro and in vivo studies, didn’t show any health hazards [6].

Sepiolite is even registered by the EU as an additive for animal feed [7, 8]. However

if there were any risks, they would be expected to be associated with inhalation

through the respiratory system. In this respect, an aspect of concern can be the

similarity of sepiolite morphology with asbestos, a notorious carcinogen. Sepiolite

particles are fibrous at a microscopic level with an average length of 1-2 µm, while

asbestos fibres have a much longer particle length, even of millimetres. Only fibres

with a length longer than 5 µm are considered a possible health hazard, although they

also need to meet other conditions, for example, biopersistence in biological tissues

for very long periods. The two minerals (asbestos and sepiolite) are also quite easy to

distinguish (i.e. XRD) and they are not usually contaminated by each other. In fact

they have a completely different geological origin. Most of sepiolite clay (including

the one employed in this thesis) has a sedimentary geological origin. They have been

formed, around 15 M years ago, by chemical precipitation in shallow lakes in periods

of arid climate when the concentration of elements (Si, Mg, Al mainly) were

suitable. These conditions are quite rare and this is one of the reasons why there are

so few commercial sepiolite deposits in the world. On the contrary, asbestos are

originated in conditions of higher pressure and temperature that produce well

crystallised and very long particles. In fact, the conditions for the formation of

sedimentary sepiolite are not compatible with the formation of asbestos and therefore

sepiolite cannot occur along with asbestos. There is other sepiolite type, very rare,

that is formed in hydrothermal conditions and whose particles have a longer length.

This sepiolite type could be contaminated with asbestos since the conditions for the

formation of this particular sepiolite are compatible with the formation of asbestos.

Sepiolite Costs

It is not easy to give an exact indication of the price of nanoclays because these

products are still mostly under development, especially if specific modifications are

required. Also, the price will definitely depend on the final customer (application

area, potential consumption, etc.). According to recent reports [9], the current price

41


of nanoclays based on organically modified montmorillonite varies between 5 and 18

€/kg and it is expected go down to 3-8 €/kg in the next decade. Conventional fillers

are much cheaper, but also vary greatly depending on the mineral and grade.

Concerning sepiolite-based products, the current price ranges between 0.7 €/Kg to 2

€/kg, depending on the type of product and whether it is modified or not. The target

price for sepiolite as a nanoclay is in the range of 3-4 €/kg [10].

2.4 References 1. R.E. Grim, Clay Mineralogy. 1968, New York: McGraw–Hill.

2. R.E. Grim, Applied Clay Mineralogy. 1962, New York: McGraw–Hill.

3. http://www.ima-mineralogy.org.

4. E. Galan, Properties and applications of palygorskite-sepiolite clays. Clay

Minerals, 1996. 31(4): p. 443-453.

5. A.J. Aznar, E. Gutierrez, P. Diaz, A. Alvarez, and G. Poncelet, Silica from

sepiolite: Preparation, textural properties, and use as support to catalysts.

Microporous Materials, 1996. 6(2): p. 105-114.

6. http://www.hse.gov.uk/lau/lacs/37-2.htm.

7. P. Suáreza, M.C. Quintana, and L. Hernández, Determination of bioavailable

fluoride from sepiolite by “in vivo” digestibility assays. Food and Chemical

Toxicology, 2008. 46(2): p. 490-493.

8. http://www.tolsa.com.

9. http://www.nanoroadmap.it.

10. Private communication with Tolsa.

42

3 Polymer / Clay Nanocomposites

3.1 Introduction

Although the intercalation chemistry of polymers mixed with appropriately modified

natural layered silicates or synthetic layered silicates has long been known [1-3],

recent findings have attracted a lot of interest and more and more research in the field

of polymer-clay nanocomposites [4-6]. Very important has been the pioneering work

carried out by Toyota Central Research Laboratories on Nylon 6-Montmorillonite

(MMT) nanocomposites. They obtained encouraging enhancements in mechanical

properties along with heat distortion temperature and decrease in permeability, with

only few weight percentage content of MMT, by an in situ-polymerisation method

[7, 8]. Melt-processing studies on polymer-clays nanocomposites followed and Vaia

and Giannelis have been among the first and most proliferous authors in the field [9].

Melt-processing is of particular interest in prospective of the industrial applications

of nanocomposites, but many problems and limitations still have to be overcome. In

fact, the simple mixture of layer-silicates with polymer melts doesn’t guarantee to

obtain nanocomposites. Actually it’s not often the case and a lot of literature

accounts for the difficulties of this task. The fundamental concept of nanocomposites

is based on the high aspect ratios and large interfaces provided by nano-fillers and

hence a substantial reinforcement obtained at small loadings, given that a perfect

dispersion of individual clays in a polymeric matrix is achieved. On the other hand,

43

CHAPTER 3 – Polymer / Clay Nanocomposites

the peculiarity of nano-fillers of having very high specific surface areas and small

dimensions simultaneously leads to a preference for agglomeration in micrometric

stacks or bundles due to Van der Waals interactions, ionic interaction and/or

hydrogen bonds. In this eventuality a traditional micro-composite would be obtained,

with poor interactions between the organic and inorganic phase. The situation is even

more challenging if the dispersing matrix is hydrophobic as it is for many important

classes of polymers. In this case the intrinsic hydrophilic nature of layered silicates

doesn’t allow a good exfoliation and dispersion because of lack of thermodynamic

driving forces. Different strategies have been pursued. A first approach takes to the

modification of clay surfaces, mainly by ion-exchange reaction with cationic

surfactants including primary, secondary, tertiary, and quaternary alkyl-ammonium

or alkyl-phosphonium cations (Fig. 3.1). These surfactants are able to intercalate into

the inter-layers galleries, swelling the clays and, at the same time, bearing a long

aliphatic tail compatible with hydrophobic polymers or, eventually, functional groups

able to react with a polymer or initiate a polymerisation reaction.

Figure 3.1. Schematic representation of modification of platelet-like nano-clays by

cation-exchange reaction with surfactant molecules. From [10].

The effect is to help the accessibility of polymer molecules in the clay inter-layers,

by increasing the clay basal distance, and the compatibility of the clays with the

polymer matrix. Because of the importance of the intercalating surfactants, a lot of

studies focused in understanding the inter-layer structure of organo-modified

44


silicates. Traditionally, based almost exclusively on WAXS analysis, the organic

chains were thought to lay either parallel to the clay, forming lateral mono- or bi-

layers, or, depending on the packing density and the chain length along with the

temperature, radiate away from the silicate surface forming extended mono or even

bimolecular tiled paraffin-like arrangement (Fig. 3.2). Such idealised structure is

based almost on all-trans conformation adopted by the alkyl chains.

Figure 3.2. Interlayer structures of intercalated surfactant molecules into the nano-

clays galleries. From [11].

Successive studies [11] based on FTIR, found the intercalated chains existing in

more complicated structures, contemplating the presence of ‘gauche’ conformations.

As the interlayer packing density or the chains length decreases (or the temperature

increases), the intercalated chains adopt a more disordered, liquid-like structure (Fig.

3.3).

Figure 3.3. Alkyl chain aggregation model. As the chain length increases the

structure changes from liquid-like to liquid crystalline-like. From [11].

45


However, the extensive use of organo-modified nano-clays in the literature has not

always given the wished results and the natural tendency to agglomeration has either

been difficult to overcome or led to thermodynamically unstable composites.

Polypropylene (PP) or polyethylene (PE)-based nanocomposites, for instance, for

which good levels of exfoliation had been claimed when processed with high shear

forces apparatus, undergo (partial) re-agglomeration of the nano-clays.

Figure 3.4. In situ SAXS-WAXS experiments during melting/recrystallisation. The

appearance of a small angle reflection after polymer melting shows partial re-

agglomeration of the clays. From [10].

Fig. 3.4 shows small angle and wide angle X-ray diffraction (SAXS and WAXS)

patterns of an isotactic polypropylene (iPP) nanocomposite during a

melting/crystallisation temperature cycle. Partial re-agglomeration is demonstrated

by the appearance of a small angle reflection when the nanocomposite melts and the

system evolves towards a thermodynamically more stable configuration.

We can understand that the ability to control the dispersion of nano-fillers in a

polymeric phase is the key issue that affects the performances of the final material

and the possibility itself to obtain nanocomposite. Depending on the thermodynamic

driving forces and the interfacial interactions between polymer matrix and nano-

clays, three main different structures can be distinguished:

46


I. Conventional micro-composites

Polymer chains are not able to enter into the interlayer space of the clays, which

preserve their agglomerated stacked structure, due to low affinity of the organic and

inorganic phases and poor interface properties. In this case we can not expect

properties far from those of conventional filled composites.

II. Intercalated nanocomposites

Polymer chains enter into the silicates interlayers and wet the clays with few polymer

layers. Hence the basal distance increase but the clays are still found in an ordered

stacked manner.

III. Exfoliated nanocomposites

Single nano-clays are completely and randomly dispersed into a continuous polymer

matrix. The distance between clays depends only from the filler loading and not from

any attractive forces.

Figure 3.5. Representation of the possible morphology arising from the interaction

of nanoclays and polymer: a) non-intercalated, conventional micro-composite; b)

intercalated nanocomposite and c) exfoliated nanocomposite. From [4].

47


3.2 Preparation of Polymer-Clay Nanocomposites

Several strategies have been considered to prepare polymer-clay nanocomposites.

We can distinguish four main groups: solution processing, in situ polymerisation,

melt processing and template synthesis.

3.2.1 Solution Processing

Solution processing is based on a solvent system in which the polymer is soluble and,

at the same time, the nano-clays are able to swell. In general, the clays are first

swollen in a solvent to form a homogeneous suspension in which the soluble polymer

is successively added. The process ends with the evaporation of the solvent or the

precipitation of the mixture, trapping the polymer chains intercalated into the inter-

layers of the clays. Polymer intercalation from solution involves a large number of

solvent molecules to be desorbed from the layered silicates in order to host polymer

chains instead. From an energetic point of view, the decrease in conformational

entropy of the confined polymer chains into the clays galleries is compensated by the

gain in translational degree of freedom during desorption of solvent molecules.

This method found extensive applications especially for water-soluble polymers.

Among the most studied there are Poly(vinyl alcohol) (PVOH) [12], Poly(ethylene

oxide) (PEO) [13], Poly(vinyl pyrrolidone) (PVP) [14], Poly(acrylic acid) (PAA)

[15]. Example of non water-soluble polymers are high density Poly(ethylene)

(HDPE) [16], with organo-modified Montmorillonite and a mixture of

Xylene/Benzonitrile as solvent, or Poly(lactide) PLA [17] and Poly(ε-caprolactone)

PCL [18] in hot Chloroform.

3.2.2 In Situ Polymerisation

In situ polymerisation involves the swelling of clays in a liquid monomer, or

monomer solution, which is successively polymerised directly in presence of

48


49

intercalated/exfoliated clays. Probably the most famous example of clay

nanocomposites, PA6-MMT hybrids obtained by Toyota research centre [7, 8], was

produced with the aforementioned method. Na+-MMT modified by �,�-aminoacid

(+H3N-(CH2)n-1-COOH, with n=2, 3, 4, 5, 6, 8, 11, 12, 18) were swollen in liquid �-

caprolactam at 100°C that was consequently ring-opening polymerised (Fig. 3.6).

Figure 3.6. Representation of the swelling of nanoclays by �-caprolactam and its

ring-opening polymerisation. From [19].

PCL-nanocomposites have also been similarly obtained. Messersmith and Giannelis

first reported on the in situ polymerisation of �-caprolactone in presence of

aminolauric modified Na-MMT at high temperature [20] and in presence of Cr3+-

exchanged fluorohectorite at 100°C for 48hrs [21]. Several studies exist on PS (and

PMMA) nanocomposites based on the free radical in situ polymerisation in presence

of organo-modified smectite nano-clays [22, 23]. A common recipe is to add organo-

modified layered clays and an appropriate initiator, i.e. N,N’-azobis(isobutyronitrile)

(AIBN), to a styrene monomer solution and carry out the polymerisation at 80 °C for

5 h. A more elegant strategy was used by Weimer et al. [24]. In this case the Na+-

Montmorillonite was cation-exchanged with a chosen ratio of “inactive”

trimethylbenzylammonium and nitroxide-bearing ammonium cation. Nitroxide

moieties are known for controlling the ‘living’ free radical polymerisation of styrene.

When the polymerisation is carried out, at 125 °C for 8 hrs, a nanocomposite with

the desired number average molecular weight, Mn, is obtained simply by tuning the

molar amount of initial monomer and nitroxide moieties. Remarkably this is among

the few examples of complete exfoliated nanocomposites.


The first to use in situ polymerisation for the preparation of PP-clay nanocomposites

was Tudor et al. [25]. The process can be schematised in a first step in which a

synthetic hectorite was treated with methylaluminoxane “MAO”, in order to remove

the acidic protons, followed by the intercalation of soluble metallocene catalyst

([Zr(η-C5H5)2Me(THF)]+), that was then able to initiate the coordination

polymerisation of PP. With a similar approach, in situ HDPE-nanocomposites were

obtained using palladium–based [26, 27] or titanium-based Ziegler-Natta catalysts

[28-30]. Next to the thermoplastic polymers aforementioned cited, in situ

polymerisation has been widely used also thermoset and elastomer/clay

nanocomposites systems. Messersmith and Giannelis [31] first reported studies of

different curing agents and curing conditions for the formation of nanocomposites

based on di-glycidyl ether of bisphenol A (DGEBA) in presence modified-

montmorillonite. A DGEBA derivative (Epon 828), cured with a polyether diamine

(Jeffamine D2000) instead formed the rubber-epoxy matrix of nanocomposite as

reported by Lan and Pinnavia [32]. The montmorillonite used were modified with n-

octylamine and n-octadecylamine and, depending on the alkyl chain length, an

intercalated (n-octyl) or exfoliated (n-octadecyl) nanocomposites were obtained.

Finally the synthesis of intercalated nanocomposites based on elastomeric

polyurethane is mentioned [33]. An organo-modified montmorillonite is swollen in a

polyol such as ethylene glycol, polyethylene glycol or glycerol propoxylate and then

cross-linked with methylene diphenyl diisocyanate.

3.2.3 Melt Processing

If the previous two methods were firstly adopted to obtain nanocomposites, melt

processing is certainly the method most interesting from an application and

economical point of view. In fact solution processing involves abundant use of

solvents, expensive and “environmentally unfriendly”, while in situ polymerisation is

often difficult to control. Melt processing involves annealing a mixture of polymer

and nanoclays, statically or under shear, above the softening temperature of the

polymer. The process is schematised in Fig. 3.7. A multitude of micro/nano-

50


composites produced by melt processing can be found in literature and most of the

main polymers have been used as dispersant phase for layered silicates. PS was first

studied [6, 30] with different types of clays and different organo-modifications, but

also PA6 [34], PP [35], EPDM [36], just to cite a few.

Figure 3.7. Schematic representation of melt-intercalation processing. From [37].

The morphologies found were claimed to be intercalated and, much less often, as

exfoliated. But, if polymer intercalation from solution processing has been justified

by a translational entropic gain of the solvent molecules balancing the polymer chain

entropic loss due to inter-layers confinement, how can we explain the same

phenomenon during melt processing? In fact, cations exchanged surfactants

molecules are tethered to the layered silicates inter-layer surfaces; hence they can not

gain any translational entropy. Moreover the silicate layers are relatively large (~ 1

µm) and unaltered during intercalation so their small translational entropy can not

contribute substantially to the hybrids formation. A first theoretical model that

explained the polymer melt intercalation in organically-modified layered silicates

was developed by Vaia and Giannelis [37, 38]. The simple mean-field lattice-based

model contemplated all the possible equilibrium states already observed

experimentally, namely immiscible, intercalated and exfoliated polymer-nanoclays.

Let’s consider an initial state in which a pair of silicate layers, distant h0 from each

other, with the inter-layer gallery completely filled by end-tethered surfactant chains,

is immersed in a polymer melt. Polymer chains will intercalate the organo-modified

clays, and then increase the inter-layer spacing up to h, if the process is

thermodynamically favourable in terms of Helmholts free energy:

51


STEhFhFF ∆−∆=−=∆ )()( 0 Equation 3.1

Composed by an internal energy contribution E∆ , associated with the establishment

of new intermolecular interactions, and an entropy contribution , associated with

the configurational changes of the various constituents. The entropic change

S∆

S∆ of

the system arises from the contributions of the confined polymer chains in the inter-

layer and the tethered chains. In fact, although the tethered chains do not gain

translational freedom, they do gain configurational freedom as the gallery height

increases.

Figure 3.8. , and as a function of the gallery height polymer

VSh∆ chainVSh∆ VSh∆ 0hh − ,

for a polymer and a layered silicate modified with octadecylammonium. Reproduced

from [37].

The overall entropy change per area of the system represented in the Fig. 3.8 might

be qualitatively divided into two regions. For a gallery height less than a critical

value hc the total entropy change is near zero. Here the penalty for polymer

confinement is compensated by the entropy gain coming from the layers separation.

For h>hc, instead, the overall entropy change is negative and the nanocomposite

52


formation is entropically unfavourable. Since the total entropy change is anyway

small, favourable interactions ( E∆ ), especially between polymer chain and silicate

layers, are expected to provide the driving force for polymer intercalation. The

interplay of entropic and energetic contributions can be interpreted by free energy

curves as in figure below.

Figure 3.9. Variation of free energy per unit area, as a function of the gallery

height , for a polymer and a layered silicate modified with octadecylammonium

group. Curves are calculated for different values of interaction energies ε

Vfh∆

0hh −

sp,sa

between polymer and silicates, supposing the interaction between polymer and

tethered chains, εsa=0. Curves I, IIa, IIb, III correspond to εsp,sa values of 0, -4, -8, -

12 mJ/m2. Reproduced from [37].

The free energy curves, shown in Fig. 3.9, can be grouped into three types. Type I is

positive for all gallery heights and hence polymer intercalation is unfavourable.

Curves of type II a-b present one or more minimum so intercalation is favourable to

certain discrete gallery heights. For type III the curve continually decreases with

gallery height, possibly leading to silicates exfoliation. However for h>h∞ (h∞ is the

53


fully extended length of the tethered chain) the entropic contribution of the tethered

chains is minimal and the penalty of polymer confinement dominates (Fig. 3.8). In

this case exfoliation requires strong favourable energetic interactions. We can

conclude that the mean-field lattice model developed by Vaia and Giannelis

describes, in first approximation, the polymer melt intercalation process in

nanocomposites. Up to a certain gallery height, the entropic penalty to confine

polymer chains can be compensated by conformational freedom of the tethered chain

gained as the layers separate. Complete layer separation (exfoliation) depends,

though, on the establishment of very favourable polymer-clay interactions to

overcome the entropy loss.

Successively, also Balazs et al. [39, 40] studied the interactions between two

surfactant-coated surfaces (layered silicate) and a surrounding polymer melt. Based

on both numerical and analytical self-consistent field (SCF) theories they tried to

draw conclusions on the equilibrium behaviour of the overall mixture. Besides

qualitatively confirming the results of Vaia and Giannelis, Balazs et al. focused on

some aspects important to design polymer-clay nanocomposites. For instance, they

revealed that significantly increasing the length of the surfactants, beyond the typical

value C12- C20, up to polymeric-like values can enhance the thermodynamic stability

of exfoliated hybrids. Moreover, studying the effect of polymer architecture on the

miscibility of polymer/clay mixtures, they found that increasing the extent of

branching at fixed molecular weight yields more miscible structures [41].

After the above brief theoretical background of the polymer melt intercalation, in the

next session we will describe some relevant examples of nanocomposites obtained by

melt processing, based on polypropylene (PP) and polyamide 6 (PA6).

PA6-Clay Nanocomposites

The preparation of nanocomposites based on PA6 has been widely investigated but

mainly limited to the in situ polymerisation method. Liu et al. [34] first melt blended

commercial PA6 with octadecylammonium-exchanged montmorillonite nanoclays in

a twin-screw extruder. WAXS analysis showed an intercalated structure for nanoclay

54


loadings above 10 wt.% while exfoliation for less than 10 wt.%. DSC tests, along

with WAXS, also stressed a strong influence of exfoliation on the crystallisation of

PA6, with heterophase nucleation effects and the appearance of γ-phase crystals in

addition to the original exclusive α-phase. Fornes et al. [42] have successively

carried out a series of investigations on the melt processing of PA6-nanocomposites.

Comparing three different PA6-matrix molecular weights (Mn=16400, 22000 and

29300 g/mol, respectively referred as LMW, MMW and HMW) they observed a

better exfoliated structure for the HMW PA6 leading to superior mechanical

properties due to the higher shear stress. Low frequency viscosity tests revealed

solid-like, non-Newtonian behaviour for HMW nanocomposites, while Newtonian

plateau for MMW and LMW. Surprisingly, though, capillary data showed lower

viscosity for HMW and MMW composites compared with pure matrix values. The

former phenomenon has been generically addressed by the authors as possibly

caused by “higher clay platelet alignment, smaller particle size and/or PA6 molecular

weight degradation”.

More recently Fornes et al. [43] compared the reinforcement effect of two different

commercial sources of clays: Kunipia-F® from Yamagata, Japan and Cloisite® Na+

from Wyoming, USA, both organo-modified with bis-(hydroxyethyl) methyl

rapeseed ammonium chloride. The first resulted in better exfoliation and better

mechanical properties, due to higher aspect ratio of single platelets, higher CEC and

higher initial basal spacing. The same authors [44] selected a series of organic amine

salts as cation exchangeable surfactants for MMT, and compared their effect on the

final properties of melt compounded PA6-nanocomposites. Three surfactant features

were identified that led to greater extent of exfoliation, higher stiffness and strength:

(1) One long alkyl tail on the ammonium ion rather than two, (2) methyl groups on

the amine rather than 2-hydroxy-ethyl-groups, (3) an equivalent rather than excess

amount surfactant on the clays.

A novel melt-compounding process was reported by Hasegawa et al. [45]. A Na+-

MMT water slurry was blended with PA6 melt in an extruder, followed by removing

of the water from a vent by vacuum (Fig. 3.10). Although WAXS and TEM revealed

55


exfoliation of the nanoclays in the PA6 matrix, the overall properties weren’t much

different from melt compounding starting from dry clays.

Figure 3.10. Representation of the “slurry” melt-compounding process. From [45].

For polymers that require high melt processing temperature, as PA6, the thermal

stability of the surfactants, usually alkyl ammonium compounds, becomes an

important issue. According to thermo-gravimetric analysis (TGA) these organic

compounds start to breakdown at temperature as low as 180°C in inert atmosphere,

and significant degradation occurs above this temperature [46, 47]. Such degradation

may affect the thermodynamics of polymer intercalation, altering the level of

exfoliation and interfacial bonding, and possibly producing side-reaction with the

polymer matrix. VanderHart et al. [48, 49] estimated, according to solid-state NMR,

that most of the quaternary ammonium on the surface of the clay is decomposed,

during PA6 melt compounding at 240 °C, as a consequence of the combination of

high temperature and mechanical shear. Surprisingly, though, extensive degradation

of the surfactant did not result in poor mixing; in fact the nanocomposites with the

best dispersion of clays also had the most extensively degraded surfactant. However

the instability of the organo-clay can adversely affect the polymer matrix itself.

Fornes et al. [50] demonstrated the degradation of PA6 matrix from a decrease in

molecular weight. The degradation was higher for a higher initial molecular weight

polymer. A significant reduction in number average molecular mass was also notices

56


by Davis et al. [51], when in situ polymerized PA6-MMT pellets were dried and

injection moulded at 300°C. The suggested degradation pathway was hydrolytic

peptide scission and/or catalytic degradation due to the presence of MMT.

These observations point to the need of more stable surfactants. In this direction goes

the research of Gilman et al. [52] that prepared PA6 nanocomposites based of MMT

modified with trialkylimidazolium cations to obtain high stability at high processing

temperatures.

PP-Clay Nanocomposites

Polypropylene has also been investigated for the preparation of nanocomposites,

especially because of the importance and wide spread use of this commodity

polymer. Differently than for PA6, the lack of polar groups in PP macromolecules

doesn’t allow direct melt intercalation into the silicates galleries; this accounts for the

difficulties in obtaining well exfoliated, PP-based nanocomposites. The general

approach has been to modify the clays surface and/or functionalise the polymeric

matrix.

Figure 3.11. Intercalation of maleic anhydride molecules into the clays galleries.

From [53].

The ability of maleic anhydride to intercalate into the clays galleries (Fig. 3.11) was

demonstrated by Tjong et al. [53]. Vermiculite clays were initially pre-treated with

hydrochloric acid solution followed by addition of maleic anhydride in presence of

acetic acid, to form organo-clays. The absence of WAXS diffraction peaks suggested

the exfoliation of the clays. Beside an effective intercalation agent for clays, maleic

57


anhydride (MA) is in general a modifying additive for the polyolefins. The melt

interaction of PP oligomers modified with either maleic anhydride (PP-MA) or

hydroxyl group (PP-OH) in octadecylammonium-exchanced MMT was first studied

by Usuki and Kato [35]. An increase in the inter-layer spacing was shown from

WAXS tests as PP-MA or PP-OH were melt compounded with organo-modified

clays. Using a higher ratio of PP-MA/clay also resulted in a better intercalation. The

authors also studied the effect of functionalisation. PP-MA with a lower maleic

anhydride content produced a reduced intercalation, demonstrating that a minimum

level of functionalisation for PP chains is necessary to intercalate silicate layers (Fig.

3.12).

Figure 3.12. Intercalation and successive exfoliation process of clays by melt-

compounding with PP-MA. From [5].

The property of PP-MA to intercalate organo-clays was successively used, by the

same research group, to obtain compatibilised three-component PP/PP-MA/MMT

nanocomposites [54, 55]. The preparation method consisted of a first step in which

the functionalised polymer (PP-MA) was intercalated into stearylammonium-

exchanged clays, followed by melt compounding with virgin PP in a twin-screw

extruder. The morphology of the nanocomposites was characterised by an

intercalated structure, when relatively high concentrations of PP-MA (~20 wt.%)

58


were used and when the compatibiliser (PP-MA) was sufficiently functionalized

(high polarity). However we have to underline that a higher functionalisation of the

compatibiliser PP-MA is not always the best solution. An excessive number of

carboxylic groups, spread on the PP oligomeric chain, can be deleterious in

decreasing the compatibility with PP matrix and inducing phase separation.

Therefore we can state that obtaining a good dispersion of organo-clays into PP

matrix compatibilised by functionalised oligomers is a compromise between high

polarity of the compatibiliser (meaning higher intercalative capacity) and

compatibility with the polymeric matrix. Manias et al. [56, 57] investigated several

other functionalisations using either random copolymers of PP (with 0.5 mol% of

functionalised comonomer) or di-block copolymers (with 1 mol% of non-PP blocks).

Different functional groups were found to promote the formation of nanocomposites

with organo-modified MMT as the common maleic anhydride but also styrene, p-

methylstyrene, hydroxyl-containing styrene and methylmethacrylate. Moreover the

same authors reported that the modification of octadecylammonium-exchanged

MMT with a semi-fluorinated alkyltrichlorosilane was directly miscible with neat

PP. Interestingly Liu and Wu [58] reported the preparation of PP-MMT

nanocomposites by a grafting/melt intercalation method. The original approach

consisted in using two intercalating agents: hexadecylammonium cations and

epoxypropylmethacrylate, an unsaturated monomer that can be tethered to the PP via

a grafting reaction. When the co-intercalated clays were blended with PP in a twin-

screw extruder an improvement in dispersion was observed because of the larger

initial inter-layer distance and stronger interactions caused by the grafting reaction.

3.2.4 Template Synthesis

A final method reported for the preparation of polymer-clay nanocomposites is the

template synthesis. The general idea is similar to the in situ polymerisation method,

with the difference that this time clays are synthesised in presence of a polymer

solution. Corrado and Xu [59] succeeded to synthesise hectorite clays in presence of

water-soluble polymers as PVP, hydroxylpropylmethylcellulose (HPMC),

59


Poly(acrylonitrile) (PAN), Poly(dimethyldiallylammonium) (PDDA), and

Polyaniline (PANI). The synthesis consists in the hydrothermal crystallisation of

silica sol, magnesium hydroxide sol, lithium fluoride and polymer in water and

reflux for 48hrs. Limitations of this method are in the incorporation of large amounts

of polymer (with an upper limit of 86 wt.% for PAN-hectorite system) and the length

of synthetic clays that, at best, is about one-third of the natural counterpart.

3.3 Properties of Polymer-Clay Nanocomposites

3.3.1 Crystallisation

The crystallinity of polymers can be dramatically modified by the presence of

layered silicates and if this is the case one has to be cautious when assessing the

correlation between microstructure and properties of polymer-clay nanocomposites.

The dispersed clay particles often act as heterogeneous nucleating agents for the

polymers, significantly reducing the spherulites size and increasing the overall

crystallinity. This effect is usually more pronounced for very low loading levels.

Above certain loadings, nanoclays may hinder the polymer mobility and then

obstacle the chains folding and the growth of well developed lamellar crystals [60].

According to the literature, enhanced crystallisation was observed for different

polymers such as PP [61, 62], PA6 [63, 64], PE [65], PET [66], etc. In some cases, a

reduction in PEO crystallisation has been reported in presence of Na+-MMT [67],

explained by the strong interactions between PEO chains and the surface of Na+,

promoting non-crystalline PEO coordination.

The effect of different particulate fillers to act as nucleating agents for polypropylene

has long been studied because of the importance of this polymer and the wide-spread

use of filled-PP for different applications. The very strong nucleating effect of talc,

for example, has been widely demonstrated [68, 69]. The influence of other fillers is

often not so clear. CaCO3, for instance, has been frequently classified as an inactive

filler. However, significant increase of nucleating effect of CaCO3 was noticed with

60


decreasing particle size and, as a consequence, increasing particle aggregation [70].

More recently, the nucleation effect of MMT in melt blended PP nanocomposites

was also studied [71-73]. Fig. 3.13 shows the reduction of spherulites size in PP after

addition of 4 wt.% of organo-modified MMT.

Figure 3.13. Polarised light micrographs of (a) neat PP (b) PP+4%clay

nanocomposites crystallised at 150°C after Yuan and Misra [74].

The heterogeneous nucleation is also explicated by a significant increase in the

crystallisation temperature, as shown in the Fig. 3.14.

Figure 3.14. DSC traces of neat PP and PP+4%clay nanocomposites. Reproduced

by Yuan & Misra [74].

61


Pozsgay et al [71] demonstrated that the ability of MMT to nucleate isotactic-PP

(iPP) depends on the organo-treatment of clay and they attributed this effect to the

alteration of interlayer MMT distance rather than to the modification of the clay

surface tension. They concluded that the nucleation occurs not on the surfaces but

rather in the interlayers of clay particles and that the nucleating effect is due to the

collapsed MMT galleries of 1nm distance. In confirmation to that hypothesis

Svoboda et al. [72] found an increase in crystallisation temperatures in PP-

MA/MMT systems containing clay tactoids but not in systems with well dispersed

MMT clays. Apart from changing the overall crystallinity and crystallisation rate, the

presence of nano-fillers can also modify the crystal structure of semi-crystalline

polymers, stabilising a metastable phase and inducing polymorphism. A typical

example is the stabilisation of γ-phase in PA6 with the addition of clays, as was

shown in early studies [34, 75]. This is important since the crystalline structure can

affect physical and mechanical properties.

Figure 3.15. Hydrogen bonding in the α and γ crystalline forms of nylon 6 as seen

from end and side-view of each crystal. Closed and open circles represent chain axes

projecting out of and into the page, respectively. From [77].

62


Some properties of PA6 based nanocomposites [34, 42, 76] may be related to such

crystalline modification rather than the effect of the nano-clay itself. Depending on

the way macromolecules are packed, PA6 can crystallise into two different crystal

forms: α-form or γ-form. In order to maximise the H-bonds in the crystal state,

polyamide chains can adopt either a full-extended or a twisted configuration. As

shown in Fig. 3.15, in the fully extended configuration (α-form), polymer chains are

oriented anti-parallel to each other, in a way that the amide linkage and methylene

units lie in the same plane and the H-bonds occur between adjacent anti-parallel

chains. This structure creates monoclinic crystals. The γ-form, instead, occurs when

H-bonds form between parallel chains and the amide linkage is twisted 60° out of

plane. This structure is described as pseudo-hexagonal. The crystal form that PA6

can assume when crystallised depends on different factors such as thermal

conditions, applied stress, moisture and additives content. For instance, it is known

[78, 79] that crystallisation at temperatures below ~130 °C favours the formation of

only γ-form, while above ~190 °C of only α-form. Both crystal forms are found for

intermediate temperatures. Fast cooling or quenching from the melt produces γ-form

[80], while successive annealing at temperatures above ~190 °C leads to the

conversion of γ-form into α-form [81]. In general we can say that fast cooling and

lower crystallisation temperatures promote the formation of γ-form, and slow cooling

and higher crystallisation temperatures leads to α-form. It seems that conditions that

limit the polymer chain mobility favour the crystallisation of PA6 in the γ-form.

Figure 3.16. TEM pictures showing the crystalline morphology of a) PA6 and b)

nanocomposites with 2.5 wt.% of clays. Reproduced by Shen et al. [82].

63


As mentioned before, the presence of additives like nanoclays can modify the crystal

structure of PA6, decreasing significantly the crystal size (Fig. 3.16) and promoting

the formation of γ-crystals (Fig. 3.17).

a) b)

Figure 3.17. WAXS spectra of PA6 nanocomposites compared with pure PA6 (a) at

the surface and (b) in the bulk of injection moulded bars. PA2 and PA3 refer to two

different organo-modified MMTs with the second highly swollen. Reproduced by

Varlot et al. [83].

A difference in structure appears between the skin and the core region of injection-

moulded specimens. All samples contain reflections at 2θ ≈ 20 ° and 23.7 °,

corresponding to the α-crystals, and at 2θ ≈ 21.3 °, corresponding to the γ-crystals,

however the skin region is much richer in γ-phase (Fig. 3.17.a) compared to the core

region (Fig. 3.17.b). That might be explained by the high shear stresses and fast

cooling imposed by the injection moulding process especially at the mould surface.

64


3.3.2 Mechanical Properties

Young’s Modulus

The addition of layered clays in a polymeric matrix has often shown remarkable

increases in Young’s Modulus. Interestingly, that increase is obtained at very low

filler content (<5 wt.%) if the clays are well exfoliated. PA6 nanocomposites

synthesised by in situ polymerisation of ε-caprolactam, show a different level of

exfoliation depending on the nature of the acid used to catalyse the reaction [84].

That is demonstrated by WAXS measurements, with an intensity peak Im that is

inversely related to the extent of exfoliation. For a decrease in Im, an increase in

Young’s modulus was observed, stressing the importance of exfoliation for the

stiffness (Table 3.1).

Table 3.1. WAXS peak intensity and Young’s modulus for nanocomposites obtained

by in situ polymerisation, catalysed with different acids. From [84].

Acid Im Young’s modulus [GPa]

Phosphoric acid 0 2.25

Hydrochloric acid 200 2.05

Isophtalic acid 255 1.74

Benzenesulfonic acid 280 1.74

Acetic acid 555 1.63

Trichloroacetic acid 585 1.67

Supposing that a good exfoliation is obtained, the size of the dispersing particle is,

obviously, of great importance. Fig. 3.18 shows the tensile modulus, measured at 120

°C, of in situ polymerised PA6 nanocomposites based on two different layered clays.

MMT (NCH in Fig. 3.18) gives a better reinforcement compared with Saponite

(NCHP in Fig. 3.18) since it is characterised by a higher aspect ratio [75].

65


Figure 3.18. Tensile modulus, measured at 120°C, of nanocomposites based on

MMT (● NCH) and Saponite (■ NCHP) clays at different loadings. From [75].

Nanocomposites produced by in situ polymerisation are sometimes difficult to

compare since the polymerisation reaction can result in matrix characterised by a

different average molecular weight and molecular weight distribution. Instead data

relative to melt compounded nanocomposites are easier to interpret. Fig. 3.19 shows

a significant and constant increase in Young’s modulus with nanoclays content up to

10 wt.%, after which the morphology is characterised by an intercalated structure

rather than perfectly exfoliated as for lower filler content.

Figure 3.19. Effect of clay content on tensile modulus of PA6-organomodified MMT

nanocomposites obtained by melt compounding. From [34].

66


A similar behaviour can be observed for PP based nanocomposites obtained by melt

compounding. In this case a better intercalation/exfoliation is enhanced by the

addition of a higher amount of PP-MA in the three components mixture PP/PP-

MA/MMT [55]. Table 3.2 shows that, as the amount of PP-MA increases also the

modulus increases. It is important to notice that there is a negligible effect of

different amounts of PP-MA on the modulus of the blend PP/PP-MA, so any

difference in the reference matrix can be excluded.

Table 3.2. Effect of PP-MA on the Young’s modulus of PP-based nanocomposites.

From [55].

Sample Filler wt. [%] PP-MA wt. [%] Young’s modulus [MPa]

PP 0 0 780

PP/PP-MA 7 0 7.2 714

PP/PP-MA 22 0 21.6 760

PPCC 6.9 0 830

PPCH 1/1 7.2 7.2 838

PPCH 1/2 7.2 14.4 964

PPCH 1/3 7.2 21.6 1010

Yield Stress and Tensile Strength

An increase in stiffness is not a sufficient proof for efficient reinforcement if not

accompanied by a simultaneous increase in yield stress or ultimate tensile stress,

since the modulus always increases when an inorganic filler is added to a polymer.

We can assume, in first approximation, that reinforcement and composite properties

are determined by factors such as matrix and filler properties, contact surface

(influenced by aspect ratio and extent of exfoliation) and strength of interaction. The

former factor is particularly important for properties of the composite like yield

stress and or ultimate tensile stress.

67


Figure 3.20. Schematic formation of hydrogen bonds in PA6/MMT nanocomposites

[5].

Without an effective interaction between matrix and reinforcing phase, with the

formation of an extended interphase region, the filler mainly acts as an inclusion,

decreasing the effective load-bearing cross-section of the matrix and, actually,

reducing the matrix strength.

Figure 3.21. Yield stress of PA6 nanocomposites for different MMT loadings. Three

molecular weights PA6 matrices are melt compounded with organo-modified MMT

clays. From [42].

PA6 has been found to be a good candidate for polymer/clay nanocomposites also

because it’s believed to form strong interactions with silicate layers via formation of

hydrogen bonds as shown in Fig. 3.20. Fig. 3.21 shows the dependence of yield

stress on MMT content and on matrix molecular weight for PA6 nanocomposites

68


obtained by melt compounding [42]. Yield stress increases with MMT content for all

formulations but more pronouncedly for high molecular weight (HMW) PA6; the

increase effect is nearly double at the highest clays content. This can be explained by

a better exfoliation for nanocomposites based on (HMW) PA6.

The results relative to the strength of PP-based nanocomposites are not always that

encouraging, mainly due to poor interfacial interactions between PP and inorganic

fillers. Referring again to the work of Hasegawa et al. [55] we can actually see a

decrease in the yield stress of the nanocomposites, compared with neat PP, for all

formulations except for PPCH 1/2 (weight ratio of nanoclays and PP-MA of 1:2)

(Table 3.3).

Table 3.3. Effect of PP-MA on the yield stress of PP-based nanocomposites [55].

Sample Filler wt. [%] PP-MA wt. [%] Yield Stress [MPa]

PP 0 0 32.5

PP/PP-MA 7 0 7.2 31.4

PP/PP-MA 22 0 21.6 32.6

PPCC 6.9 0 31.9

PPCH 1/1 7.2 7.2 29.5

PPCH 1/2 7.2 14.4 34.6

PPCH 1/3 7.2 21.6 31.7

The co-intercalated PP/MMT nanocomposites studied by Liu and Wu [58] (refer to

Section 3.2.3.2) show a better reinforcement effect. The tensile strength considerably

increases with the MMT content up to a value of 5 wt.%, after which the trend

becomes less pronounced (Fig. 3.22). The strong interfacial interactions shown by

these nanocomposites are probably due to the grafting reaction between the PP

chains and the co-intercalated unsaturated monomer epoxypropylmethacrylate.

69


Figure 3.22. Effect of clay content on tensile strength of co-intercalated PP/MMT

nanocomposites. From [58].

Strain at Break

The effect of nanofillers on the property of the strain at break has also been studied.

The usual effect for thermoplastic matrices, such as PMMA [85], PS [86], PA6 and

PP is a considerable reduction in the strain at break. Table 3.4 shows the strain at

break for different PP/MMT nanocomposites.

Table 3.4. Effect of PP-MA on the strain at break of PP-based nanocomposites.

From [55].

Sample Filler wt. [%] PP-MA wt. [%] Strain at break [%]

PP 0 0 >150

PP/PP-MA 7 0 7.2 >150

PP/PP-MA 22 0 21.6 40.3

PPCC 6.9 0 105

PPCH 1/1 7.2 7.2 7.5

PPCH 1/2 7.2 14.4 8.6

PPCH 1/3 7.2 21.6 5.6

70


71

There is a dramatic decrease from values higher than 150 %, for neat PP, down to 5-

8 % for PP/PP-MA/MMT nanocomposites. However, it has to be stressed that the

presence of PP-MA could have an important effect and needs to be better

understood. In fact, if we consider the polymer blend PP/PP-MA, with 22 wt.% of

PP-MA (third entry of Table 3.4), the strain at break is already reduced to 40.3 %.

Fornes et al. [42] reported the properties at failure of PA6/MMT nanocomposites, for

three different molecular weight PA6 and two crosshead speeds. As shown in Fig.

3.23, the strain at break is reduced with MMT content, in particular for the lower

molecular weight matrix (LMW) that, we remember, was characterised by a coarser

dispersion of MMT. The high molecular weight matrix (HMW) seams to preserve

ductility to a certain extent.

a)

b)

Figure 3.23. Elongation at break of three MW PA6 nanocomposites for different

MMT loadings, tested at crosshead speed of (a) 0.51 cm/min and (b) 5.1 cm/min.

From [42].

3.3.3 Barrier Properties

The high aspect ratio of layered silicates has been found to significantly reduce the

gas permeability in exfoliated nanocomposites films, by the creation of a “tortuous

path” (Fig. 3.24) that reduce the diffusiveness of gas molecules.


Figure 3.24. Formation of a “tortuous path” in polymer-clay nanocomposites.

We can understand that the plate-like shape of the nano-clays is particularly efficient

in maximising the path length that a diffusing molecule must travel, because of the

high aspect ratio compared with others fillers shapes (i.e. sphere or cube). According

to the argument developed by Nielsen [87] the tortuousity factor τ is defined as the

ratio of the actual distance d’, that a diffusient must travel, to the shortest distance d,

that it would travel in absence of obstacles. τ can be expressed in terms of the length

L, width W, and volume fraction φ of the sheet-like particle as follows:

ϕτWL

dd

21'+== Equation 3.2

The effect of tortuosity on the permeability is expressed as:

τϕ−

=1

P

S

PP

Equation 3.3

where PS and PP represent, respectively, the permeability of the nanocomposite and

the pure polymer. Supposing that the layered clays are arranged perpendicular to the

direction of diffusion, a key role is played by the aspect ratio, as showed in the Fig.

3.25.

72


Figure 3.25. Relative permeability plotted as a function of aggregate width, for

different sheet lengths. Effect of exfoliation. From [88].

This simple model has been usefully employed by different authors to interpret

experimental data. Lan et al. [89] measured the permeability to carbon dioxide of

partially-exfoliated polyimide-based nanocomposites. The curve fitting was obtained

for aspect ratio of 192 (Fig. 3.26), lower than the value of 1000 expected for well

exfoliated MMT. Interestingly, Yano et al. [90] found a ten-fold decrease in water

vapour permeability for polyimide with the addition of only 2wt.% of nanoclays.

73


Figure 3.26. Relative permeability for different loading of MMT clays. The best

fitting is for aspect ratio 192, much lower than 2000 expected for completely

exfoliated MMT platelets. From [89].

3.3.4 Fire Retardancy

The addition of few weight percentages of nanoclays has showed to be beneficial for

the fire retardancy of different polymer matrices.

Figure 3.27. Heat release rate of PA6 and PA6/silicate nanocomposites (5 wt.%).

Reproduced from [91].

74


Cone calorimeter is the main bench-scale method to evaluate important parameters in

the fire retardant behaviour of a material such as the heat release rate (HRR), the

peak of HRR, smoke production and CO2 yield. In a typical cone calorimeter

experiment, a sample is exposed to a given heat flux and the heat release rate and the

mass loss rate are recorded as function of time. Gilman et al. first reported detailed

investigations on the flame retardant properties of PA6/layered clays, followed by

studies on nanocomposites based on other polymer matrices [91-93].

a) b)

c)

Figure 3.28. Residues of combustion of: (a) EVA with 5 phr organoclays; (b) EVA

with 5 phr MWCNTs; (c) EVA with 2.5 phr pure MWCNTs and 2.5 phr organoclays.

From [94].

As typical example, Fig. 3.27 shows the HRR plot of PA6 and PA6 nanocomposites

(5 wt.% of MMT), with a reduction of 63 % in HRR peak. The suggested mechanism

of flame retardancy in nanocomposites arises from the formation of a char layer,

obtained through the collapse of intercalated/exfoliated structures, that act as a

barrier both to mass and energy transport. As the fraction of clays increases, the

75


amount of char increases and the rate at which heat is released decreases. The

morphology of nanocomposites again plays an important role. Good nanoclay

dispersion is certainly precondition for fire retardancy but has been reported that an

intercalated structure can be more effective than a completely exfoliated structure.

Usually also the simple observation of the residues of combustions of a material

gives useful information. From Fig. 4.28 we can see three different compounds of

EVA filled with clays (Fig.4.28.a), MWNTs (Fig.4.28.b) or both the fillers

(Fig.4.28.c), after a cone calorimeter test [94]. The use of both clays and MWNTs

seems to work in synergism leaving a very thick and homogeneous char without

presence of cracks, preventing then by the emission of volatiles.

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85

4 Needle-like Clay Nanocomposites

4.1 Introduction - The Importance of Shape

Most of the research reported in literature on polymer/clay nanocomposites has been

focusing on platelet-like clays, in general smectite clays such as MMT. The

exfoliation of layered silicates and preparation of homogeneous nanocomposites is

seriously limited by the strong tendency of platelet-like to agglomerate due to their

extended contact surface.

Figure 4.1. Surface area to volume ratio (A/V) as a function of the aspect ratio (l/d)

of cylindrical particles. Reproduced from [1].

86

CHAPTER 4 – Needle-like Clay Nanocomposites

Instead, the dispersion of inorganic fibres in a nanometre scale is a much easier

challenge due to a relative small contact surface. This is confirmed by the graph in

Fig. 4.1, that shows the surface area to volume ratio A/V in function of aspect ratio

(l/d) of a cylindrical particle, for a given particle volume. Values of l/d<0 correspond

to platelet-like particles, while l/d>0 correspond to rod-like particles. We can see that

A/V increases faster for platelet-like than for rod-like particles with respect to the

aspect ratio [1]. Hence, for an equivalent volume of particle, the platelet-like have a

higher contact surface. Furthermore, the mechanical reinforcement potential of fibres

is higher than platelet, as has been theoretically described by Gusev [2] and Van Es

[3], for unidirectional composites. Fig. 4.2, for instance, shows the graph of relative

Young’s modulus for unidirectional composites filled with platelets and fibres of

different aspect ratios, according to the Halpin-Tsai and Mori-Tanaka models. The

two models are equivalent for a fitted value of the shape factor ζ (ζ=2/3 l/t for flakes

or platelets and ζ=(0.5 l/t)1.8 for fibres) of the Halpin-Tsai equations (see Page 122-

123).

Figure 4.2. Reinforcement effect of platelets and fibres in unidirectional composites,

for different aspect ratios, according to Halpin-Tsai and Mori-Tanaka models.

Reproduced from [3].

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As we can see, the fibres offer a better reinforcement and approach the theoretical

maximum, defined by the simple rule of mixture, faster and for lower values of

aspect ratio. As described in Chapter 3, layer silicates are not only found in the

common platelet-like (or flake-like) shape but also in a fibre-like shape, as apply for

clay minerals such as sepiolite, attapulgite, and palygorskite. The objective of this

chapter is an overview of the literature relative to fibre-like (or needle-like) clay

nanocomposites.

4.2 Preparation methods

The preparation of needle-like clays nanocomposites has followed strategies similar

to the most common platelet-like clays nanocomposites, and can be classified into

the same categories.

4.2.1 In Situ Polymerisation

The in situ polymerisation method, as already stated earlier, involves the dispersion

of the filler into a monomer solution followed by polymerisation of the reaction

mixture.

a) b)

Figure 4.3. SEM micrographs of PA6/attapulgite nanocomposites obtained by in

situ polymerisation. The pictures refer to filler concentration of: a) 2 wt.% and b) 5

wt.% (right). From [4].

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The classical PA6 in situ polymerisation method has been recently adopted by

different authors [4-6]. Shen et al. [4] prepared PA6/attapulgite nanocomposites by

in situ polymerisation of the needle-like clays in ε-caprolactam monomer.

Attapulgite was previously pre-modified with cetyltrimethylamonium bromide

(CTAB) and reacted with toluene-2,4-diisocianate (TDI). SEM investigations were

carried out, on formic acid-etched surfaces, to evaluate the dispersion of attapulgite

in PA6 matrix, as shown in Fig. 4.3.

In situ polymerisation was also employed by Ozdilek et al. [6] in order to prepared

PA6/boehmite nanocomposites by hydrolytic polymerisation of ε-caprolactam in

presence of an aqueous boehmite suspension. Boehmite is not classified as a clay

mineral but it is included in this overview for the similarities with the other needle-

like clays. Boehmite, also known as γ-AlOOH, is a needle-like colloidal particle with

sizes in the nanometre scale and high anisotropy, hydrothermally synthesised from

aluminium alkoxide precursors. Fig. 4.4 shows the well defined shape of the

boehmite particles produced.

Figure 4.4. TEM picture of boehmite stabilised in n-propanol. Reproduced from [6].

The in situ polymerisation results into a homogeneous nanocomposite with well

dispersed needle-like particles, as it appears from the TEM pictures in Fig. 4.5. For

filler content below 7.5 wt.%, the boehmite particles are randomly dispersed into

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PA6 matrix, whereas at 9 wt.% they seem to show parallel nematic orientation. The

same authors, in a successive publication [5], described the in situ polymerisation of

PA6 in presence of Titanate-modified boehmite.

Figure 4.5. TEM pictures of PA6/boehmite nanocomposites, obtained by in situ

polymerisation, referring to (a) 7.5 wt.% and (b) 9 wt.% of filler. From [6].

With this functionalisation it was possible to obtain a more stable dispersion in ε-

caprolactam and a higher content of the filler without using excessive amount of

water.

Figure 4.6. TEM pictures of two different concentrations of Ti-modified oehmite/PA6

nanocomposites: 7 wt.% (left) and 15 wt.% (right). From [5].

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More in details, a dispersion of Ti-modified boehmite needles in n-propanol was

combined with ε-caprolactam to form a homogeneous mixture. When n-propanol

was completely extracted in rotavapor, the solid mixture monomer/filler left was

used for the polymerisation reaction. Morphological investigations by TEM (Fig.

4.6) show that Ti-modified boehmites are homogeneously distributed into PA6

matrix. Some regions of the higher concentrated nanocomposites (15 wt.%) seem to

show a nematic orientation of the needles.

In situ polymerisation was also used, by Du et al. [7], in order to prepare PE-based

nanocomposites. Palygorskite needle-like clays were first activated with the initiator

TiCl4, adsorbed on the clay surface, and then introduced, along with AlR4, hexane

and ethylene under pressure, into a reactor where the polymerisation took place.

Polyethylene chains were generated simultaneously on different activated points of

the palygorskite surface, resulting into a highly entangled “macromolecular comb”.

The so-produced PE/palygorskite nanocomposites were successively melt-blended

with a commercial grade PE and then processed in a twin screw extruder at 220 °C.

Epoxy nanocomposites based on sepiolite clays were reported by Zheng [8].

Different amounts of sepiolite were dispersed in diglycidyl ether of bisphenol A

(DGEBA) at 130 °C, followed by the addition of an hardener (epoxy:hardener =

1:0.875).

Figure 4.7. TEM picture of sepiolite in epoxy matrix. From [8].

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The cured epoxy nanocomposites presented a filler dispersion described by the

authors as “tree branched” (Fig. 4.7). More recently, Xue and Pinnavia [9]

systematically studied the effect of palygorskite functionalisation on the preparation

and properties of epoxy-based nanocomposites. The needle-like clays were

stoichiometrically silylated, assuming a density of four silanol sites per square

nanometre of palygorskite, with different reagents such as: γ-

aminopropyltrimethoxysilane (APTMS), N-dodecyltriethoxysilane (DTES) and

1,1,1,3,3,3-hexamethyldisalazane (HMSZ). Modified and not-modified clays were

successively dispersed in DGEBA and cured in order to obtain rubbery or glassy

epoxy nanocomposites. If pristine clays were not well dispersed in epoxy pre-

polymer and tend to sediment before curing, the silylation modification substantially

improved the clay dispersability.

The preparation of elastomer/sepiolite nanocomposites has been reported by

Bokobza et al. [10].

Figure 4.8. TEM pictures of elastomer/sepiolite nanocomposites containing 5phr of

clays. From [10].

Clay particles were dispersed into the monomer, 2-hydroxyethyl acrylate, with a high

shear mixing apparatus for several hours and the mixture was polymerised, under UV

illumination, after addition of 0.05 wt.% of a photo-initiator (Irgacure 819, from

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CIBA). The TEM picture in Fig. 4.8, relative to the elastomer/sepiolite

nanocomposite containing 5 phr of clays, shows good sepiolite dispersion.

4.2.2 Melt Compounding

One of the first needle-like particles nanocomposites prepared by melt-compounding

was reported by Souma [11-13]. The author obtained the composite by mixing

Xonotlite with unplasticized PVC in a heating roll. Xonotlite is a needle shaped

calcium silicate and it is synthesized by hydrothermal reaction of Ca(OH)2 and SiO2.

Two techniques were employed to promote the disaggregation of Xonotlite particles.

One was a mechanical atomisation by jet air mill (abbreviated as X(JA)) and other

was the addition of hexamethyl disiloxane (HMDS) directly into the reaction system

(abbreviated as X(HMDS)). In a series of successive publications, Acosta et al. [14-

16] focused on the use of sepiolite as filler for polypropylene. Sepiolite was used

pristine, as received, or previously modified (esterified) with different aliphatic

organic acids, by means of a condensation reaction of the superficial hydroxyl groups

with the acid groups of the reagents. Three sepiolite concentrations (10, 25, 40 wt.%)

were obtained melt-compounding treated and untreated-sepiolite with PP in a

Brabender Plasticorder internal mixer, preheated to 200 °C. The rotor speed was set

at 60 rpm and the mixing time at 15 min. The same authors presented the production

of a hybrid composite in which different amounts of sepiolite were added to a

PP/glass fibre composite [17]. The approach was intended to examine any synergistic

effects of the two fillers and also to partially replace the glass fibres for a cheaper

material, with obvious economic benefits. The binary systems PP/glass fibre (glass

fibre content: 10, 20, 30 wt.%) and PP/sepiolite (sepiolite content: 10, 25, 49 wt.%)

were melt-compounded in a roll mixer at 190 °C, in the needed amounts in order to

obtain ratios of 10/20 and 20/10 wt.% of glass fibre/sepiolite in the final hybrid

composites.

Recently Wang and Sheng [18, 19] reported on the preparation of polypropylene/org-

attapulgite nanocomposites. The needle-like clays were organo-modified by first

reacting with γ-methacryloxypropyl trimethoxysilane followed by a grafting of butyl

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acrylate. Varying amount of org-attapulgite clays were then melt-blended with PP in

a mixer apparatus at a temperature of 185 °C and a residence time of 10 min at 32

rpm, to obtain nanocomposites with filler concentration from 1 to 7 wt.%. Fig. 4.9

refers to a 5 wt.% PP/org-attapulgite.

Figure 4.9. TEM picture of PP/org-attapulgite nanocomposites with 5 wt.% filler

content [19]. Although no scale bar appears in the micrograph, the diameter of the

attapulgite is estimated at 40nm.

4.3 Properties of Needle-like Clay Nanocomposites

4.3.1 Crystallisation

It is well known that the presence of inorganic particles can dramatically change the

crystallisation behaviour of polymers. In order to understand the relation between

structure and physical and mechanical properties of nanocomposites materials it is

extremely important to evaluate the crystallinity and the crystallisation.

Acosta et al. [16] first studied the crystallisation kinetic of PP composites containing

Sepiolite surface-treated with isobutyric acid. The authors conducted both isothermal

and non-isothermal kinetics studies. For the isothermal crystallisations the samples

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were heated to 473 K, held at that temperature for 5 min, in order to minimize the

effect of thermal history, and successively cooled down and held at three different

crystallisation temperatures (399 K, 401 K and 403 K); the results were interpreted

with the Avrami’s equation [20-22], which describes the crystallisation kinetics and

can be written as:

)()())1ln(ln( tnLnKLnX t +=−− Equation 4.1

Where Xt is the progressive volume fraction transformed into crystal until the time t,

K a scaling factor and n the Avrami constant. The constant n gives insight into the

nucleation process (homogeneous/heterogeneous) and the shape of the growing

crystal (rod/disc/sphere/sheaf).

Non-isothermal crystallisations were carried out at three different cooling rates (10, 5

and 2.5 K/min) and studied in terms of the Harnish and Muschik’s method. The

above kinetics studies concluded that sepiolite treated with isobutyric acid acted as

nucleating agent for polypropylene, increasing both crystallinity ratio and rate

(evaluated by Avrami’s values K). Moreover the needle-like clays induced the

formation of a second ordered PP structure, claimed to derive from an interphase on

the particle periphery and demonstrated by the graphic representation of the

Avrami’s equation, which showed two well-defined slopes. Wang and Sheng [23,

24] studied the isothermal and non-isothermal crystallisation of PP/org-attapulgite

nanocomposites prepared by melt compounding and already described in session

4.2.2. Four crystallisation temperatures (121, 122, 123 and 124 °C) were chosen for

the isothermal crystallisation tests and the results explained in terms of the Avrami’s

equation and Hoffman’s theory (Fig. 4.10). Org-attapulgite particles incorporated in

PP matrix acted as heterogeneous nuclei, dramatically increasing the crystallisation

rate, with values of Avrami’s exponent n increased from 2.39-2.81 to 4.24-4.93.

Only minor effects were found on the equilibrium melting temperature and degree of

crystallinity. Polarised optical microscopy observations revealed that the spherulites

size decrease and the number increase with attapulgite content.

95


Figure 4.10. Avrami’s plot for (a) neat PP and PP/ATP nanocomposites: (b)

PP/ATP 1 wt.%, (c) PP/ATP 3 wt.% and (c) PP/ATP 5 wt.%, at three different

crystallisation temperatures [24].


The mechanical properties of polymeric materials have been shown to be remarkably

improved when nanocomposites are formed with low filler content. Wang and Sheng

[19] have recently investigated the tensile properties of PP/attapulgite produced by

melt-compounding. Both pristine and organo-modified attapulgite clays were used as

nano-fillers in PP matrix. The Young’s modulus and yield stress of nanocomposites

with different filler content is presented in Fig. 4.11. All samples show an increased

modulus, though the enhancement of the org-attapulgite is much larger than that of

pristine attapulgite.

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a) b)

Figure 4.11. Young’s modulus and yield stress for PP/Org-ATP nanocomposites in

function of the filler content [19].

The presence of any amount of non-modified attapulgite decrease the yield stress of

nanocomposites compared to the neat PP, meaning weak interfacial interactions

matrix/filler. The yield stress of PP/org-attapulgite nanocomposites, however, shows

a maximum increase of 13 %, compared to neat PP, relative to a 3 wt.% content of

clays. Unfortunately the properties at break are not presented by the authors and can

not be commented on. The properties of PE/palygorskite nanocomposites prepared

by Du et al. [7] via an in situ polymerisation method are shown in the Table 4.1.

Table 4.1. Mechanical properties of in situ polymerised PE/palygorskite

nanocomposites [7].

Sample Filler wt.

[%]

Tensile strength

[MPa]

Elongation at break

[%]

Impact strength

[kJ/m2]

PE 0 31.4 180.2 50.7

PB1 1 38.1 172.1 82.9

PB2 1.7 32.3 170.2 70.9

PB3 2.2 26.1 166.7 60.2

PB4 3.3 22.4 160.8 42.8

Interestingly, the tensile strength is significantly enhanced with just 1 wt.% of

palygorskite needle-like clays without the trade-off of a reduced ductility. In fact the

97


elongation at break for that concentration was only limitedly reduced while the

impact strength remarkably increased. Xue and Pinnavia [9] have studied the

mechanical properties of epoxy polymers reinforced by different modified-

palygorskite and the results are presented in Table 4.2.

Table 4.2. Mechanical properties of rubbery epoxy/palygorskite nanocomposites [9].

Nano-particle and

nano-particle loading [%]

Tensile strength

[MPa]

Tensile modulus

[MPa]

Tensile elongation

[%]

None 0.62 (6.8 %) 2.8 (5.0%) 25.0 (8.0%)

Palygorskite

2 0.71 (8.2 %) 4.4 (7.7 %) 25.3 (9.2 %)

5 0.74 (8.5 %) 4.0 (6.9 %) 20.7 (8.9 %)

10 1.32 (7.3 %) 5.0 (6.4 %) 28.1 (9.0 %)

APTMS-Palygorskite

2 0.81 (4.7 %) 4.7 (7.0 %) 24.0 (10.3 %)

5 0.89 (5.5 %) 3.5 (6.0 %) 28.4 (8.8 %)

10 1.03 (9.2 %) 4.4 (7.3 %) 21.8 (9.8 %)

DTES-Palygorskite

2 1.18 (2.3 %) 4.7 (5.2 %) 30.9 (2.6 %)

5 0.98 (4.2 %) 3.5 (9.1 %) 30.8 (5.5 %)

10 1.47 (6.3 %) 6.2 (8.4 %) 26.0 (7.7 %)

HMSZ-Palygorskite

2 0.97 (5.0 %) 4.8 (3.8 %) 24.3 (8.1 %)

5 1.00 (7.2 %) 5.4 (6.0 %) 21.1 (6.0 %)

10 1.35 (6.4 %) 6.7 (9.5 %) 21.2 (13.4 %)

Organo-MMT/Epoxy

0 ~0.5 ~2.6 -

2 ~0.8 ~5.5 -

5 ~1.5 ~7.5 -

10 ~3.3 ~13 -

Numbers in parentheses are the relative standard deviation.

98


The as-received palygorskite provide some improvements in the modulus of epoxy

matrix but little improvements in tensile strength at filler loading of 2 wt.% and 5

wt.%. The abrupt enhancement at 10 wt.% is explained by an increase in the

viscosity of the pre-polymer mixtures that prevents the settling-out of the filler. The

organo-modification of the palygorskite provides a better dispersion and hence better

mechanical properties, especially at low filler content (2 wt.% and 5 wt.%).

Comparing the reinforcing effect of palygorskite and organo-modified-MTT we can

observe that the second provide better mechanical properties for higher loading of

filler. However the level of organo-modification for palygorskite is <0.2 % of the

level used to compatibilise MMT, with obvious cost benefits.

4.3.3 Rheology

The rheology of particulate suspensions is governed by factors such as the structure,

size and shape of the dispersed phase as well as its orientation distribution and the

strength of interaction between the dispersed and dispersant phase. Understanding

the rheological properties of polymer/clay nanocomposites is crucial in gaining

insights into the processability and the structure-properties relations of these

materials. Shen et al. [4] performed rheological studies on PA6/attapulgite

nanocomposites and investigated the effect of varying loadings of the fibrous silicate

on the linear and non-linear viscoelastic properties. The results were interpreted in

terms of solid-like behaviour and percolated network structure and comparisons were

drawn between the behaviour of needle-like and platelet-like clays. The linear

dynamic viscoelastic curves for PA6 and PA6/attapulgite nanocomposites are shown

in Fig. 4.12. At the temperature of 160 °C and the frequency range of 10-1-102 rad s-1,

PA6 exhibits the usual behaviour of homopolymer melts with low Mw distribution,

i.e. and , where G’ and G’’ and ω are respectively the storage shear

modulus, the loss shear modulus and the angular frequency [25]. The addition of

clays significantly modifies the viscoelastic response. At high frequencies the

behaviour of G’ and G’’ is qualitatively similar while at low frequency range

2' ω∝G ω∝''G

99


PA6/attapulgite nanocomposites have higher G’, G’’ and viscosity when compared

with PA6 and show a monotonic increase with clay content. Moreover all the

nanocomposites display significantly reduced frequency dependence. Especially at

low frequency range, G’ and G’’ becomes almost independent from the frequency,

which is characteristic of materials with solid-like behaviour, i.e.

and .

0' ω∝G0'' ω∝G

Figure 4.12. Small amplitude strain sweep (A) and frequency sweep (B-D) at 260 °C

for neat PA6 (a) and PA6 nanocomposites with attapulgite content of: (b) 2 wt.%, (c)

3 wt.%, (d) 4 wt.% and (e) 5 wt.%. From [4].

The above frequency dependence of G’ and G’’ is similar to that observed by

Krishnamoorti and Giannelis [26] for PCL-MMT nanocomposites, explained by the

formation of a percolated structure of the dispersed nanoclays in the polymer matrix

(Fig. 4.13).

100


Figure 4.13. Schematic description of the polymer/needle-like clay percolating

structure. From [26].

Simplifying the lattice model in Isichenco’s review [27], Shen et al. formulated a

‘grafting-percolating’ model able to describe the percolation thresholds of attapulgite

sticks (aspect ratio l/d=30-40) in a polymer matrix.

Figure 4.14. 2D sketch of the percolation lattice model at Φ<Φc (A), Φ=Φc (B) and

Φ>Φc. Black occupied lattices represent the sticks percolated, while the grey

occupied lattices represent the sticks unpercolated. The symbol ‘X’ indicates the

lattices occupied by grafted polymer chains. From [4].

101


Applying a Monte Carlo simulation to the 3D lattice model (Fig. 4.14), the authors

were able to predict a percolation occurring at a volume fraction of attapulgite Φc of

0.02, or, alternatively, a weight fraction of 3-4 wt.%.

4.4 References

1. H. Fischer, Polymer nanocomposites: from fundamental research to specific

applications. Materials Science & Engineering C-Biomimetic and

Supramolecular Systems, 2003. 23(6-8): p. 763-772.

2. A.A. Gusev, Numerical identification of the potential of whisker- and platelet-

filled polymers. Macromolecules, 2001. 34(9): p. 3081-3093.

3. M. Van Es, Polymer-Clay nanocomposites-the importance of particle

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104

PART 2: Experimental Results and

Discussion

105

5 The Potential of Sepiolite Nanoclay

as a Reinforcement for Polymer

Composites

5.1 Introduction

Numerous publications are already available on polymer/clay nanocomposites (as

shown in Chapter 3 and 4), though most of it deals with smectite clays (i.e. MMT).

The main originality of this PhD thesis, with respect to the existing literature, is to

study sepiolite clay as nanofiller for polymer nanocomposites and to draw

comparisons with the commonly studied smectite clays (Chapter 5, 6, 8). This

comparison is interesting because of the intrinsic shape difference between the two

clays (needle-like shape for sepiolite and platelet-like shape for smectite clays),

which is expected to be particularly important in oriented structures. In fact, few are

the works on oriented polymer/clay nanocomposites and Chapter 7 represents an

important contribution to this research area. Another open question in the scientific

community is the effect of polymer/clay interphase in the final nanocomposite

mechanical properties. Chapter 9 will try to give an answer to this, by preparing

model nanocomposites via in-situ polymerisation.

106

CHAPTER 5 – Sepiolite Nanoclay as a Reinforcement for Polymer Composites

The following chapter is meant to provide fundamental information about sepiolite

nanoclay. It is reasonable to believe that each nano-filler behaves differently and also

that the same nano-filler can perform differently in different polymers, according to

its specific characteristics. First, the shape and dimensions of the filler are studied via

scanning and transmission electron microscopy, followed by Brunauer-Emmett-

Teller (B.E.T.) analysis in order to determine the specific surface area. The

dehydration of such nanoclay is also quantified by thermo gravimetric analysis in

inert atmosphere. For the first time in the literature, the stiffness of sepiolite nanoclay

is directly measured by nano-bending tests using scanning probe microscopy (SPM).

The last paragraph of this chapter will predict the reinforcing efficiency of sepiolite

nanoclay in two thermoplastic polymers (PP and PA6) which will be used in the

proceeding of this thesis, using micromechanical theories. A particular attention will

be given to the comparison between sepiolite and more commonly used smectite

clays (i.e. Montmorillonite (MMT)) in terms of theoretical reinforcement.

5.2 Experimental

5.2.1 Materials

The sepiolite clays (Pangel) were supplied by Tolsa (Spain), and used as received.

The porous substrate used for the SPM nano-bending tests, in a holographic epoxy

membrane of 1µm pore size, provided by Eindhoven Technical University (The

Netherlands) [1].

5.2.2 Characterisation Techniques Morphological Analysis

Morphological studies were carried out using a Scanning Electron Microscope

(SEM), Jeol JSM-6300F, and Transmission Electron Microscope (TEM), Jeol JEM

107


2010 TEM. The samples were prepared by deposition of a sepiolite water suspension

on a metal substrate, followed by gold coating (for the SEM), or directly on carbon

coated TEM copper grids (Agar Scientific).

BET

The equipment used for the specific surface area characterisation is a Micromeritics

Gemini 2360, working in multipoint BET mode. Powder samples were dried over

night in an oven at 80 °C, followed by a treatment (degassing) at 200 °C under

vacuum for 2 hours prior to analysis, using VacPrep060. The tests are performed at

-197 °C (in a liquid nitrogen bath).

TGA

Thermo Gravimetric Analyses (TGA) were performed with a TA instrument Q500

on about 10mg samples. The tests were performed at a scanning rate of 20 K/min up

to 1000 °C, under inert atmosphere (N2) conditions.

Scanning Probe Microscopy

The instrument used is an NTegra (NT-MDT, Russia) operating with a closed-loop

system to minimize piezo drift during measurements. Force-displacement curves and

images were obtained using standard SPM probes (Nanosensors, Switzerland, spring

constant ~ 2 Nm-1).

5.3 Results and Discussion

5.3.1 Morphological Analysis Sepiolite appears as a very fine powder at naked eyes. Only an investigation at

electron microscope can reveal us more about the nature of the constituent

elementary particles. Fig. 5.1 shows an SEM micrograph of sepiolite clays deposited

on a porous substrate.

108


Figure 5.1. SEM micrograph of sepiolite clays dispersed on a porous substrate.

Some agglomerates of several microns can be observed, constituted by smaller

needle-like shaped particles. The fibrous nature of this clay can be easier

distinguished in the TEM picture in Fig. 5.2.

Figure 5.2. TEM micrograph of sepiolite on carbon-coated copper grids.

The sample can be considered to be very homogeneous although few impurities are

present. These appear as darker spots of irregular plate-like shape, probably been

different smectite clays.

109


At higher magnifications (Fig. 5.3a-b) we can notice that individual needles are quite

regular in diameter but have a wider range of lengths.

a)

b)

Figure 5.3. TEM micrographs of sepiolite clays on carbon coated TEM grids.

Distribution of: a) lengths and b) diameters. Black arrows underline single fibres

measurements.

Applying a statistical approach on a significant number of pictures, the dimensions of

the needles can be derived after post-processing with Image Pro® software (Fig.5.4)

110


a)

500 1000 1500 20000

10

20

30

40

50

60

70

Freq

uenc

y

Sepiolite Length [nm]

b)

10 15 20 25 30 35 40 45 500

10

20

30

40

50

Freq

uenc

y

Sepiolite Diameter (nm)

Figure 5.4. Distributions of: a) sepiolite lengths and b) sepiolite diameters.

111


The length of the nano-fibres varies between 200nm and 2µm, with an average of

about 650nm while the diameter varies between 15nm and 35nm, with an average of

24nm. The nanofiller average aspect ratio s (L/D) is then 27 but it ranges between 10

and 130.

In comparison, montmorillonite is a platelet-like clay with about 1nm thickness and

up to few microns width. The aspect ratio should be higher than 1000 but it is usually

in the range of 100-1000 because stacks of few intercalated plates are often found

rather than individual particles.

5.3.2 B.E.T. Measurements

The acronym B.E.T. (derived from the authors of the equation Brunauer, Emmett and

Teller) describes a technique that allows determining the specific surface area of a

solid sample [2]. Such parameter, usually expressed in m2/g, is extremely important

in fields like catalysis, for instance, where the reactivity depends on the surface area

of a given catalyst. In this study, it is important to know the surface area of sepiolite

to estimate the extension of the interface matrix/filler in polymer/clay

nanocomposite.

The determination of the specific surface area by B.E.T is based on the adsorption of

a gas on to the surface of a solid sample. An adsorptive (nitrogen in this case) is

admitted to the solid in controlled increments. After each dose of adsorptive, the

pressure is allowed to equilibrate and the quantity of gas adsorbed is calculated. The

gas volume adsorbed at each pressure (and at one constant temperature) defines an

adsorption isotherm. The specific surface area is determined by the quantity of gas

required to form a monolayer over the external surface of the sample, which can be

calculated by fitting the adsorption isotherms with the BET equation [3]. The BET

model describes reasonably well adsorption for relative pressures P/P0 between 0.05

and 0.35, and then the fitting is limited to such interval of pressure. In Table 5.1, the

values relative to a series of isotherms are reported.

112


Table 5.1. Volume of N2 gas adsorbed on sepiolite clay at different pressures.

Relative Pressure, P/P0 Pressure [mmHg] Vol. adsorbed [cm3g-1 at STP]

0.0999 77.57 86.424

0.1498 116.39 89.980

0.1998 155.21 93.173

0.2498 194.05 96.246

0.2998 232.85 99.379

The BET equation assumes the form:

admon

admon

ado

CVCPP

CVPPVPP )1(/1

)/1(/ 0

0

−+=

− Equation 5.1

Where C is a constant and is the volume of gas required to get one complete

monolayer. If we plot

admonV

)/1(/

0PPVPP

ado

− versus , a straight line can be fitted, with

slope and intercept .

oPP /


monCV/1

From the slope and intercept C and can be determined. From the linear fitting of

the BET plot in Fig. 5.5, the values obtained are C=-52.884 and =67.452. The

specific surface area can then be simply calculated by

admonV

admonV

g

Aad

montotalBET V

NVS

σ=, Equation 5.2

Where N is Avogadro’s number (6.022E+23), σA is the cross sectional area of N2

molecules (0.162 nm2) and Vg is molar volume of adsorbed gas (17.3 cm3mol-1) at

STP. It can then be calculated that the specific surface area of sepiolite clay is 300

m2g-1, in good accordance with the literature [4-7].

113


0.10 0.15 0.20 0.25 0.300.0010

0.0015

0.0020

0.0025

0.0030

0.0035

0.0040

0.0045

(P/P

0)/(V

ad*(

1-P/

P 0))

P/P0

Figure 5.5. B.E.T. plot of )/1()/( 0PPVPP ad

o − versus relative pressure C

and can be calculated from the linear fit of the data point, since the slope is

and intercept is .

)/( oPP .

admonV


monCV/1

The specific surface area calculated by BET method does not correspond simply to

the total external surface of the clay. In fact, sorption in sepiolite occurs also in the

internal channels. Hence, from the adsorption isotherms it is measured the sum of the

internal and external surface area; the inner surface area accounting for an important

fraction of the total (40 %) [4].

For comparison, the specific surface area of MMT clay calculated with the same

technique and reported in literature is much lower [8-10]. It usually varies between

25 m2g-1 and 40 m2g-1, for pristine clay and it is further reduced after surface

functionalisation.

This value is surprisingly low. The specific surface area can be extrapolated from the

one of a single platelet. A value of around 1000 m2g-1 can be calculated, taking a

density of 2 g/cm3 and approximating the particle to a rectangular prism of

114


1µmx1µmx1nm. The same calculations applied to the sepiolite, considering a

cylinder of 24 nm diameter and 650 nm length, give a specific surface of 100 m2g-1,

which is not far from the experimental value once the contribution of the internal

porosity is excluded. The low specific surface area measured for MMT can be

attributed to the denser structure, in which stacks of platelets are collapsed on each

other, effectively reducing the total accessible surfaces. In fact, it is believed that

BET measures only the external specific surface area, while the total specific surface

area, in the case of completely exfoliated clays, has been quantified in 658 m2g-1,

from TEM micrographs analysis [11]. The observation above is believed to be an

indication of the difficulties in separating individual platelet-like clays (i.e. MMT)

rather than fibre-like clays (i.e. sepiolite). Therefore sepiolite should be

comparatively easier to disperse.

5.3.3 Thermal Properties

As already mentioned one the most important characteristics of sepiolite clay,

already exploited in industrial applications, is their excellent adsorptive properties.

This is mainly due to the high specific surface area, high porosity and surface activity

of such clay.

In this section we are interested in studying the desorption of molecules, like water,

inevitably present on sepiolite. The dehydratation process of sepiolite was studied by

thermogravimetric analysis (TGA), in which the clays are subjected to a temperature

scan up to 1000 °C, at the rate of 20 K/min, in inhert atmosphere, with a flow rate of

nitrogen gas of 40 ml/min. Fig. 5.6 clearly shows four main weight losses attributed

respectively, for ascending temperature, to the release of adsorbed and zeolitic water,

the release of the first structural water, the release of the second structural water and

the dehydroxylation of the Mg-OH groups (Table 5.2) [12].

115


0 200 400 600 800 1000

80

85

90

95

100

3.3%

2.6%

3.5%

10.3%

Wei

ght [

%]

Temperature [°C]

Figure 5.6. Dehydration of sepiolite clays under temperature scan. The dashed lines

represent the temperature window at which clays are typically subjected during

composites preparation (extrusion and compression moulding).

The water adsorbed on external surfaces and the zeolitic water from the nanoporous

tunnels is removed at relatively low temperatures. The elimination of coordinated,

structural water starts when the zeolitic water is lost and ends when dehydroxylation

begins. Folding of the sepiolite crystals occurs when some structural water has been

removed. This process, reversible for temperature below 350 °C, becomes

irreversible once all the structural water molecules are removed and partial

dehydroxylation has occurred, forming an anhydride form. Finally, the remaining

Mg-OH hydroxyl groups are released at ~850 °C.

116


Table 5.2. Weight losses of sepiolite clays.

1st weight loss 2nd weight loss 3rd weight loss 4th weight loss

Amount 10.3 % 3.5 % 2.6 % 3.3 %

Physical

phenomenon

Adsorbed and

zeolitic water

First structural

water

Second structural

water

dehydroxylation

of the Mg-OH

It is important to notice, at this stage, that the temperature range at which the clays

will be subjected to for all the future nanocomposites preparation is between 200 °C

to 260 °C. At these temperatures the dehydration is completely reversible and there

are no structural changes in the sepiolite nanoclay [13].

5.3.4 Mechanical Properties and Nano-Bending Tests

One of the prerequisites for using a nano-filler as a reinforcement for polymeric

materials is the necessity of a good mechanical performance of the filler itself. Direct

experimental observations of the mechanical properties of individual nanoclays are

non trivial due to the difficulties in performing nanometre-scale tests and, to the best

of the author’s knowledge, computer simulations [14] and bulk measurements [15]

are the sole approaches being used so far to estimate such information.

The stiffness of nanoclays is assigned by some authors to be around 170 GPa, relying

on the properties of a perfect mica crystal, whose crystal structure resembles that of

MMT [16, 17]. Theoretical calculations contemplate even higher values [18].

In this section, a direct measurement of the mechanical properties of single sepiolite

nanoclays is performed via scanning probe microscopy (SPM). SPM has been chosen

for these experiments because of the possibility of controlling very accurately the

deflection of a sample, via the z-piezo displacement, as a function of an applied

force. A schematic of the experimental test, which can be described as nano-bending,

in analogy with similar works in the literature [19-24], is presented in Figure 5.7.

117


Figure 5.7. Schematic representation of SPM bending test of sepiolite nano-fibres

suspended on porous substrate.

Sepiolite nano-fibres, lying over the substrate pores, are first located using the SPM

in tapping-mode. The SPM cantilever is then brought into contact with the nano-fibre

at the midpoint along its suspended length, the operating mode is switched from

tapping to contact mode and finally a force-deflection (F-D) curve is obtained setting

an appropriate z-piezo scanner range (i.e. 50 nm down and back 200 nm up). An x-y

closed-loop feedback piezo-scanner is employed in order to have an accurate and

reliable positioning of the cantilever.

a) b)

-50 0 50 100 150 200-2.0

-1.8

-1.6

-1.4

-1.2

-1.0

Def

lect

ion

[nA]

Heigth [nm]

Figure 5.8. SPM tests: a) image of a sepiolite nanoclay laying on the substrate and

b) typical force-displacement curve.

118


Fig. 5.8 shows the SEM image of sepiolite clays laying on the substrate and a typical

F-D curve relative to the bending test performed on the fibre itself.

The slope of the curve is directly related to the elastic modulus of the nano-clay.

Supposing it is a simply supported beam and behaves within the limits of small

deformation theory

δIFLE48

2

= Equation 5.3

Where F is the force applied, L the free length of the sepiolite, δ the deflection at the

midpoint of the length and I the moment of inertia, which for a circular cross section

is

64

4DI π= Equation 5.4

Where D is the diameter of the cross section. The bending test results of four

different sepiolite fibres are presented in Table 1.

Table 5.3. Results of the nano-bending tests.

Fibre Diameter [nm] Length [nm] Elastic modulus [GPa]

1 50 1155 171

2 50 900 215

3 80 900 166

4 80 900 267

The importance of accurate measurements of the fibres dimensions, in particular of

the cross-section area should be stressed. In fact, small variations in fibre dimensions

give rise to a significant change in the values of stiffness. The values of elastic

moduli reported in Table 5.3 are affected by an uncertainty of about 20 % and more

tests are needed to make the results statistically more solid.

119


5.3.5 Micromechanical Models

This chapter has dealt so far with gaining basic knowledge on sepiolite nanoclay

fibres. The motivation was to understand and predict the properties that such

nanofillers can impart to polymer matrices, when forming a nanocomposite.

Traditional micromechanical models can provide substantial understanding toward

this purpose, and in particular the theory of short fibre reinforced composites.

Shear Lag Model

Among the most widely used short fibre composite models is the shear lag model in

which a cylindrical inclusion (fibre), surrounded by a matrix, is oriented parallel to

the external load direction. In such system the tensile stress is transferred from the

matrix to the fibre by interfacial shear stresses, supposing a perfect filler/matrix

interface (Fig. 5.9).

a) b)

Figure 5.9. Schematic illustration of the concept of the shear lag model: a)

unstressed and stressed system and b) variation of the shear stress and strain in the

matrix as a function of the radial position. From [25].

120


The variation of the axial stress and the interfacial shear stress along the fibre length,

, can be expressed as [25] x

⎥⎦

⎤⎢⎣

⎡⎟⎠⎞

⎜⎝⎛−= )(sec)/(cosh1 nshrnx

rnxE ff εσ Equation 5.5

and

( )nshr

nxEnfi secsinh

2⎟⎠⎞

⎜⎝⎛=

ετ Equation 5.6

where is the fibre elastic modulus, fE ε is the strain of the matrix (which is

supposed to be equal to the overall strain of the composite), is the axial distance

from the fibre midpoint,

x

r is the fibre radius, is the fibre aspect ratio (L/D) and

is a dimensionless constant given by

s n

⎥⎥⎦

⎤

⎢⎢⎣

⎡

+=

)/1ln()1(2

fEE

nmf

m

ν Equation 5.7

where mν is the Poisson’s ratio of the matrix and is the volume fraction of filler in

the composite. If equations 5.5 and 5.6 are plotted as a function of , it can be

observed that the tensile stress is zero at the fibre ends and assumes a maximum

value in the centre. Vice versa the shear stress at the interface is a maximum at the

ends and zero in the centre. If the fibre is long enough, the tensile stress can build up

until a plateau value,

f

x

plateaufσ , at which the strain in the fibre is equal to the strain in

the matrix and approximately equal to the composite itself (isostrain condition). The

longer the fibre the higher is the fraction of the applied load that the fibre can

withstand. The reinforcing efficiency, instead, decreases as the fibre length is

reduced since the contribution of the parts of the fibre which are not fully loaded (the

121


fibre ends) are proportionally more significant. If the fibre is too short the stress

doesn’t build up to the plateau value and it shows a maximum that is smaller than

plateaufσ .

a)

b)

c)

Figure 5.10. Predicted variations in a) fibre tensile stress and c) interfacial shear

stress along the length of a glass fibre (schematically represented in b)), in

polyester/30 % glass fibre composite, subject to an axial tensile strain of 10-3, for

two fibre aspect ratios. Redrawn from [25].

122


It comes natural to define a critical length, , which is the minimum length for

which the maximum of tensile stress at the midpoint of the fibre closely approaches

the value of

cL

plateaufσ . This critical length depends on the elastic constants of fibre

and matrix but also from the diameter of the fibre itself, and that is why it is better to

define a critical aspect ratio . It can be demonstrated [25] that the critical aspect

ratio is approximately equal to

cs

nsc

3≈ Equation 5.8

with , defined in Equation 5.7, which becomes bigger when the fibre volume

fraction increases and the ratio of fibre and matrix elastic constants decreases.

n

In figure 5.9, the critical aspect ratios of fibres, of the same elastic constant as

sepiolite, in two polymer matrices (PP and PA6) are plotted as function of the filler

vol.%. For the calculation of the following values have been used: n

fE = 200 GPa; = 1.35 GPa; = 3 GPa; PPmE

6PAmEPPmν = 0.35 and

6PAmν = 0.40.

The dashed horizontal lines in Fig. 5.11 represent the distribution of aspect ratios

measured from TEM micrographs in Section 5.3.1. The average aspect ratio line

( =27) lies above the curve representing the critical aspect ratios for PA6/Sep

composite for different filler vol.%, which guarantees an efficient reinforcement of

sepiolite nanoclay in PA6, according to the shear-lag model. The same is not valid

for softer matrices as PP. In fact the theoretical aspect ratio required for PP/Sep is

always higher than the average sepiolite aspect ratio, in the filler vol.% range

investigated. This means that, in average, the filler will only carry a smaller fraction

of the load applied to the composite. Nonetheless, as seen before (5.3.1), the

distribution of dimensions of the nanofiller is quite broad especially for what

concerns the length, which can assume values up to few microns. Hence, the aspect

cs

123


ratio can be over 100. In conclusion, for PP/Sep composite, only the higher end of

the filler aspect ratio distribution is able to efficiently reinforce the polymer matrix.

0 2 4 6 8 100

102030405060708090

100110120130

PA6/Sep

PP/Sep

Crit

ical

asp

ect r

atio

, sc

Filler Vol. [%]

Figure 5.11. Critical aspect ratios as a function of the sepiolite vol.% for two

polymer composites (solid lines). The dashed horizontal lines represent the average

aspect ratio (middle) and the lowest and highest values of the aspect ratio

distribution (bottom and top), measured from TEM micrographs.

Halpin-Tsai Model

Another very popular composite theory is expressed by the Halpin-Tsai equations.

The advantage of such model is the possibility of predicting with a simple expression

all the elastic constants of composite materials as a function of the aspect ratio of the

filler, along, obviously, with the constituent properties and the volume fractions of

the two phases. The model will also be employed to compare the theoretical

reinforcement of fibre-like and platelet-like fillers, in order to relate sepiolite with the

more commonly used smectite clays (i.e. Montmorillonite).

The Halpin-Tsai equation can be written as

124


f

f

m

c

EE

ηϕζηϕ−

+=

11

Equation 5.9

where:

⎟⎟⎠

⎞⎜⎜⎝

⎛+

⎟⎟⎠

⎞⎜⎜⎝

⎛−

=

ζη

m

f

m

f

EEEE

1 Equation 5.10

In this formula the following parameters can be identified:

Ec = composite elastic modulus

Em = matrix modulus

Ef = filler modulus

φf = volume fraction of the filler

ζ = shape factor

It should be stressed that the Halpin-Tsai set of equations was initially developed for

(and it is usually used with) fibre reinforced composites. However it can be adapted

to different filler shapes by choosing a suitable shape factor ζ. The parameter ζ can

be found by fitting experimental data. In this study, though, it will be taken from the

scientific literature.

Van Es [16], in his PhD thesis, proposed different shape factors for fibre-like and

platelet-like reinforcements (Table 5.4), by comparing the Halpin-Tsai equation with

the Mori-Tanaka theory. The last assumes spheroidal filler, which can then

intrinsically contemplate different shapes, by changing the dimensions along the

three principal axes.

125


Table 5.4. Shape factors ζ, for fibre-like and platelet-like reinforcement, relative to

the elastic moduli E11, E22, E33, along the principal directions 1, 2, 3 as in Fig. 5.12.

From [16].

Fibre Platelet

E11 ζ=(0.5s)1.8 ζ=2/3s

E22 ζ=2 ζ=2/3s

E33 ζ=2 ζ=2

The parameter s, defined already in the previous section, is the aspect ratio of the

filler. It is generally defined as the ratio of the longest and shorter dimension. It will

be then the ratio of the length over the diameter, for fibres, and the ratio of width

over the thickness for platelets. The principal directions 1, 2 and 3 of the composite

are related to ones of the two oriented fillers, are exemplified in Fig. 5.12.

3311

22

Figure 5.12. Principal directions of composites relative to the ones of oriented

fillers.

For unidirectionally oriented material (in the direction of the fibre length or in the

plane of the platelet), it can be stated that

E= E11 Equation 5.11

126


The calculation of the moduli of a composite with randomly oriented fillers, instead,

would need the application of orientation-averaging mathematical transformations to

the unidirectional case, which is rather cumbersome to use. However, an

approximated solution, inspired by the laminate theory [25], can instead be used.

For a 2D random isotropic material, the modulus of the composite can be written as

)(625.0)(375.0 //2 ⊥− ⋅+⋅≈ EEE randomD Equation 5.12

For a random 3D isotropic material, according to the shape of the filler, the modulus

of the composite can be written as [16]

)(816.0)(184.0 //3 ⊥−− ⋅+⋅≈ EEE FibrerandomD Equation 5.13

or

)(51.0)(49.0 //3 ⊥−− ⋅+⋅≈ EEE PlateletrandomD Equation 5.14

//E and are the elastic moduli of the composite in the longitudinal direction (for

the fibre is E

⊥E

11 and for the platelet is E11 or E22) and in the transverse direction (E33)

respectively and they coincide with highest and lowest value.

Assuming that the elastic modulus of sepiolite is about 200 GPa (Paragraph 5.3.4),

the elastic modulus of PP and PA6 are respectively 1.3 GPa and 3 GPa and the filler

volume fraction is fixed at 5 vol.%, the plots in figure 5.13 and 5.14 can be

generated. Figure 5.13 compares the reinforcing effect of fibre-like and platelet-like

fillers, unidirectionally oriented and of different aspect ratios, in two polymer

matrices: PP and PA6.

127


1 10 100 1000 10000 100000

1

2

3

4

5

6

7

8

9

10

rule of mixtures

rule of mixtures

Fibre Platelet

PA6

PP

E c/E

m

Aspect ratio, s

Figure 5.13. Reinforcement of 5 vol.% of fibre-like and platelet-like filler,

unidirectionally oriented (1D) in two polymer composites. The dashed vertical line

shows the average sepiolite aspect ratio. Arrows show the smallest aspect ratio

necessary to reach theoretical reinforcement (rule of mixtures).

Both fillers, for aspect ratios sufficiently high, have the same ultimate reinforcement,

reaching the same plateau which corresponds to the rule of mixtures (dashed lines).

The plateau relative to PP based composites is higher than the one relative to PA6

based composites for the simple reason that for the same filler the reinforcing effect

is proportionally more for a softer matrix than for a stiffer matrix (PP has lower

elastic modulus than PA6).

An interesting feature lies in the difference between the two filler shapes for the

same matrix. Fibre-like fillers approach maximum reinforcement already for aspect

ratios of 100, while platelet-like fillers need aspect ratios higher than 2000. It can be

concluded then that, in the unidirectional case, fibres are more effective than

128


platelets. This is however different for the situation of randomly distributed fillers

(Fig. 5.12).

1 10 100 1000 10000 100000

1.0

1.5

2.0

2.5

3.0

3.5

4.0

4.5

5.0

Fibre Platelet

E c/Em

Aspect ratio, s

Figure 5.14. Reinforcement of 5 vol.% of 3D randomly oriented fibre-like and

platelet-like fillers in PP matrix. The dashed line shows the average sepiolite aspect

ratio. Arrows show the smallest aspect ratio necessary to reach theoretical

reinforcement (rule of mixtures).

Figure 5.14 shows the reinforcement of 3D randomly oriented fibre-like and platelet-

like fillers, of different aspect ratios, in PP matrix. The observation that fibre fillers

reach the maximum reinforcement for aspect ratios much smaller than those

necessary in the case of platelet filler still holds. This effect is also more prominent

since randomly distributed platelets need an aspect ratio of 10000. However the

plateau relative to platelet fillers is twice as high as for fibres filler. It can therefore

be concluded that, in the case of randomly oriented filler, platelets are more effective

than fibres.

129


5.4 Conclusions

This chapter gives an overview of some important properties of sepiolite, the filler

which will be used throughout the rest of this thesis. Sepiolite clay is shown to be

composed of elementary particles, which peculiar shape resembles that of rigid

needles. Average dimensions are 650 nm in length and 24 nm in diameter, with a

mean aspect ratio (L/D) of 27, but ranging between 10 and 130. The specific surface

area of such clay reaches 300 m2g-1, which is partially due to the internal porosity

created by channels running through the entire length of the nano-needle. Sepiolite

has high adsorptive capacity and a particular great affinity to water molecules.

Adsorbed water normally accounts for 10-15 % of the total weight of the clay. Upon

heating to temperatures of about 300 °C, such water can be reversibly desorbed. At

temperatures higher than 350 °C, instead, also the structural water is removed and

sepiolite undergoes irreversible dehydroxylation and crystal folding. For the first

time in literature, the mechanical properties of nanoclays have been evaluated via

nano-bending tests on individual sepiolite particles. The elastic modulus of sepiolite

was found to be about 200 GPa. Finally the reinforcement of sepiolite, in two

thermoplastic matrices (PP and PA6), is predicted by means of the shear-lag model

and the Halpin-Tsai equations. Sepiolite was also compared with more frequently

used smectite clays (i.e. montmorillonite), which showed that sepiolite is expected to

give a more effective reinforcement in unidirectionally oriented composites. Vice

versa platelet-like clays should provide a higher reinforcement in the case of 3D

random composites.

5.5 References

1. J.C.A. Van der Werf, Monodisperse Holographic Membranes, in Chemistry

and Chemical Engineering. 2006, Eindhoven Technical University: Eindhoven,

The Netherlands.

130


2. S. Brunauer, P.H. Emmett, and E. Teller, Absorption of gases and

multimolecular layers. Journal of Americal Chemical Society, 1938. 60: p. 309.

3. H.J. Butt, K. Graf, and M. Kappl, Physics and Chemistry of Interfaces. 2003,

Weinheim: Wiley-VCH.

4. A.J. Aznar, E. Gutierrez, P. Diaz, A. Alvarez, and G. Poncelet, Silica from

sepiolite: Preparation, textural properties, and use as support to catalysts.

Microporous Materials, 1996. 6(2): p. 105-114.

5. M. Radojevic, V. Jovic, and D. Vitorovic, Study of sepiolite from Goles

(Kosovo, Yugoslavia). I. Sorption capacity. Journal of the Serbian Chemical

Society, 2002. 67(7): p. 489-497.

6. J.L. Bonilla, J.D.D. Lopezgonzalez, A. Ramirezsaenz, F. Rodriguezreinoso,

and C. Valenzuelacalahorro, Activation of a Sepiolite with Dilute-Solutions of

Hno3 and Subsequent Heat-Treatments .2. Determination of Surface Acid

Centers. Clay Minerals, 1981. 16(2): p. 173-179.

7. J.D.L. Gonzalez, A.R. Saenz, F.R. Reinoso, C.V. Calahorro, and L.Z. Herrera,

Activation of a Sepiolite with Diluted Hno3 Solutions Followed by Thermic

Treatment - Study of the Specific Surface. Clay Minerals, 1981. 16(1): p. 103-

113.

8. P. Praus, M. Turicova, S. Studentova, and M. Rits, Study of

cetyltrimethylammonium and cetylpyridinium adsorption on montmorillonite.

Journal of Colloid and Interface Science, 2006. 304: p. 29-36.

9. S. Miao, Z. Liu, B. Han, H. Yang, Z. Miao, and Z. Sun, Synthesis and

characterization of ZnS-montmorillonite nanocomposites and their application

for degrading eosin B. Journal of Colloid and Interface Science, 2006. 301: p.

116–122.

10. K. Boukerma, J.Y. Piquemal, M.M. Chehimi, M. Mravčáková, M. Omastová,

and P. Beaunier, Synthesis and interfacial properties of

montmorillonite/polypyrrole nanocomposites Polymer, 2006. 47(2): p. 569-576.

11. B. Chen and J.R.G. Evans, Preferential intercalation of clay in polymer-clay

nanocomposites. Journal of Physical Chemistry B, 2004. 108: p. 14986-14990.

131


12. W.X. Kuang, G.A. Facey, and C. Detellier, Dehydration and rehydration of

palygorskite and the influence of water on the nanopores. Clays and Clay

Minerals, 2004. 52(5): p. 635-642.

13. E. Galan, Properties and applications of palygorskite-sepiolite clays. Clay

Minerals, 1996. 31(4): p. 443-453.

14. S. Grigoras, A.A. Gusev, S. Santos, and U.W. Suter, Evaluation of the elastic

constants of nanoparticles from atomistic simulations. Polymer, 2002. 43(2): p.

489-494.

15. A. Kelly, Strong Solids. 1973, Oxford: Clarendon Press.

16. M. Van Es, Polymer-Clay nanocomposites-the importance of particle

dimensions. 2001, TU Delft: Delft.

17. D. Shia, C.Y. Hui, S.D. Burnside, and E.P. Giannelis, An interface model for

the prediction of Young's modulus of layered silicate-elastomer

nanocomposites. Polymer Composites, 1998. 19(5): p. 608-617.

18. O.L. Manevitch and G.C. Rutledge, Elastic properties of a single lamella of

montmorillonite by molecular dynamics simulation. Journal of Physical

Chemistry B, 2004. 108(4): p. 1428-1435.

19. A.H. Barber, S.R. Cohen, and H.D. Wagner, Measurement of carbon nanotube-

polymer interfacial strength. Applied Physics Letters, 2003. 82(23): p. 4140-

4142.

20. A.H. Barber, S.R. Cohen, A. Eitan, L.S. Schadler, and H.D. Wagner, Fracture

transitions at a carbon-nanotube/polymer interface. Advanced Materials, 2006.

18(1): p. 83-87.

21. Q.H. Xiong, N. Duarte, S. Tadigadapa, and P.C. Eklund, Force-deflection

spectroscopy: A new method to determine the Young's modulus of

nanofilaments. Nano Letters, 2006. 6(9): p. 1904-1909.

22. E.W. Wong, P.E. Sheehan, and C.M. Lieber, Nanobeam mechanics: Elasticity,

strength, and toughness of nanorods and nanotubes. Science, 1997. 277(5334):

p. 1971-1975.

23. J.P. Salvetat, A.J. Kulik, J.M. Bonard, G.A.D. Briggs, T. Stockli, K. Metenier,

S. Bonnamy, F. Beguin, N.A. Burnham, and L. Forro, Elastic modulus of

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ordered and disordered multiwalled carbon nanotubes. Advanced Materials,

1999. 11(2): p. 161-165.

24. J.P. Salvetat, G.A.D. Briggs, J.M. Bonard, R.R. Bacsa, A.J. Kulik, T. Stockli,

N.A. Burnham, and L. Forro, Elastic and shear moduli of single-walled carbon

nanotube ropes. Physical Review Letters, 1999. 82(5): p. 944-947.

25. D. Hull and T.W. Clyne, An introduction to composite materials. Second ed.

1996, Cambridge: Cambridge University Press.

133

6 Polypropylene / Sepiolite

Nanocomposites

6.1 Introduction

Polypropylene/sepiolite nanocomposites were prepared by melt compounding in a

mini-extruder apparatus. The dispersion of sepiolite in PP has been improved either

by the use of surface modification of the clay (silanisation) or by use of

functionalised polymers. The often used maleic anhydride modified polypropylene

(PP-g-MA) is compared with a custom-made functionalised polymer, PP-acid, in

respect to the filler dispersion and filler reinforcement efficiency. For that purpose,

morphological and mechanical studies are carried out by means of SEM, TEM and

mechanical tensile tests. In addition, the nanocomposites are characterised by WAXS

and DSC techniques in order to assess the effect of the nanofiller on the crystalline

structure of the PP matrix. The use of PP-PEO and the clay surface modification

resulted in a better nanofiller dispersion compared to traditional PP-g-MA modified

systems. Sepiolite acts as nucleating agent for the crystallisation of PP and seems to

lead to an orientation of the α-phase crystals.

134

CHAPTER 6 – PP / Sepiolite Nanocomposites

6.2 Experimental

6.2.1 Materials

The i-PP used was Moplen® HP500H with melt flow index MFI 1.8 g/10min (at 230

°C/2.16 Kg). The pristine sepiolite Pangel® were supplied by Tolsa (Spain), and

used as received or surface modified with alkyl-silane at TNO (The Netherland). The

two fillers above will be referred to as Sep and Sep-sil, respectively, in the

proceedings of this chapter. The PP-g-MA used is a commercial grade VINBOND®

P series (VB100; 1% grafted MA; MFI 6.0 g/min at 230 °C/2.16 Kg; Mw=151

kg/mol) from Vin Enterprise Ltd. PP-acid is a customer-made product from Baker

Petrolite. This is a carboxylic acid terminated PP, of Mn=1200 gmol-1 and with an

acid number of 30.

6.2.2 Nanocomposites Preparation

Nanocomposites were prepared by a two-steps blending process in a mini twin-screw

extruder DSM Micro 15 at 200 °C for 10 min at 200 rpm. First, sepiolite clays (Sep

or Sep-sil) were mixed with the functionalised polymer (1:1 weight ratio) and/or PP

homopolymer to make a masterbatch at 10 wt.% of filler, which was subsequently

diluted with neat PP homopolymer to obtain nanofiller concentrations of 1 wt.%, 2.5

wt.% and 5 wt.%. The composition of materials studied is listed in Table 6.1.

135


Table 6.1. Compositions of PP/sepiolite nanocomposites.

Sepiolite

(Sep or Sep-sil)

wt. [%]

PP-acid

wt. [%]

PP-g-MA

wt. [%]

PP+1%Sep 1 - -

PP+2.5%Sep 2.5 - -

PP+5%Sep 5 - -

PP+PP-acid+1%Sep 1 1 -

PP+PP-acid+2.5%Sep 2.5 2.5 -

PP+PP-acid+5%Sep 5 5 -

PP+PP-g-MA+1%Sep 1 - 1

PP+PP-g-MA+2.5%Sep 2.5 - 2.5

PP+PP-g-MA+5%Sep 5 - 5

PP+1%Sep-sil 1 - -

PP+2.5%Sep-sil 2.5 - -

PP+5%Sep-sil 5 - -

6.2.3 Nanocomposites Characterisation

Morphological Analysis

Morphological studies were carried out using Scanning Electron Microscopy (SEM)

analysis (Jeol JSM-6300F) on gold-coated, cold fractured surfaces and Transmission

Electron Microscopy (TEM) analysis (Jeol JEM 2010 TEM) on ultra-thin samples

obtained by a microtome.

X-ray

Wide Angle X-ray Scattering (WAXS) spectra were recorded with a Siemens

Diffractometer D5000, where the X-ray beam was Ni-filtered CuKα (λ=1.5405 Ǻ)

and radiation operated at 40 kV with a filament current of 40 mA. Corresponding

136


data were collected from 5 to 30 ° at a scanning rate of 0.01 °/min. The samples

analyzed were films of thickness 100 µm, hot pressed at 220 °C for 10 min.

DSC

Non-isothermal crystallization analyses were performed with a Differential Scanning

Calorimeter (DSC) TA Q1000. All samples (2.0±0.1 mg) were firstly heated to 220

°C and kept at that temperature for 5 min to remove any thermal history and then

cooled at a rate of 10 K/min.

Tensile Tests

Tensile tests were conducted in a universal testing machine (Instron 5584), equipped

with a 1 kN load cell, standard grips and Merlin software, according to the standard

ASTM D-638. The test specimens were dog-bone shaped with a length of 60 mm and

a thickness of 1 mm, according to the type V dimensions indicated by the same

standard. Specimens were obtained by compression moulding at 220 °C for 10min.

Rheology

The rheology measurements were performed on a controlled strain rheometer

(Advanced Rheometer, AR 2000) equipped with an environmental chamber and

parallel plates geometry (25 mm diameter). The specimens were disk-shaped (1.5

mm thickness and 25 mm diameter) and were tested at 200 °C. The measurements

were conducted over a frequency range of 0.01–100 Hz at the shear strain of 0.5 %,

within the linear viscoelastic region.

TGA

The thermal stability of PP/sepiolite nanocomposites was estimated by thermo

gravimetric analysis on a TA Q500. About 10 mg of sample was cut off from tensile

test specimens and tested both in air and nitrogen atmosphere. After equilibration at

30 °C, the samples were heated up to 1000 °C at a rate of 20 K/min.

137



6.3.1 Morphological Analysis

The dispersion of the nanofiller in the polymer matrix is a crucial aspect for the

performances of nanocomposites. SEM pictures of cold fractured surfaces of

different nanocomposite systems are shown in Fig. 6.1 and will be discussed in

particular concerning the influence of the different functionalised polymers and filler

surface modification on the sepiolite dispersion. Fig. 6.1.a shows a micrograph of

PP+Sep nanocomposite with filler concentration of 2.5 wt.%. The system is

characterised by micrometer sized clusters of sepiolite which will act as inclusions in

a nearly filler-free polymeric matrix. Only minor improvements in dispersion can be

observed in case PP-g-MA is employed as surface-active polymer (Fig. 6.1.b).

Sepiolite nanofibres are again mainly found in the form of clusters, which are

inhomogeneously distributed within the polymer matrix. Instead the morphology of

nanocomposites shown in Fig. 6.1.c appears significantly different. In this system the

use of PP-PEO leads to much finer filler dispersion in the polymeric matrix, where

aggregates of sepiolite are no longer evident. The silane coated sepiolite (Sep-sil),

without the use of any functionalised polymers in the blend, shows an intermediate

behaviour. Sep-sil at low filler content (1 wt.%) can be nicely dispersed in PP (Fig.

6.1.d), while at concentration of 5 wt.% is already found agglomerated.

138


a) b)

c) d)

e)

Figure 6.1. SEM micrographs of: a) PP+2.5%Sep; b) PP+PP-g-MA+2.5%Sep; c)

PP+PP-acid+2.5%Sep; d) PP+1%Sep-sil and e) PP+5%Sep-sil. White circles

underline sepiolite clusters. A significant improvement in the dispersion of sepiolite

in PP matrix is evident with the use of PP-PEO and Sep-sil, where no agglomerates

of nanoclay are found in nanocomposites at 2.5 wt.% filler load.

The morphology of the nanocomposites has also been investigated by TEM. This

analysis, besides confirming the SEM observations, underlines a characteristic

139


fracturing of the sepiolite fibres as a consequence of the mechanical stress during the

compounding process.

a) b)

Figure 6.2. TEM picture of: a) sepiolite dispersion on TEM grids b) sepiolite in PP

matrix after compounding. A reduction in fibre length is evident in the processed

nanocomposites as a consequence of melt blending in mini-extruder.

As can be seen in Fig. 6.2.a-b, a reduction of fibre length is evident in the processed

nanocomposites compared with the original sepiolite. The length of the fibres is

reduced although a quantification of such reduction is difficult to make. In fact

Fig.6.2.a is a 2D picture (sepiolite dispersed on TEM grids) and shows the real

dimensions, while Fig.6.2.b is a slice of a 3D composite in which sepiolite fibres are

randomly distributed in all directions and therefore the length observed is only a

projection on the plane of the microtome cut.

6.3.2 Crystal Structure and Crystallisation Behaviour

Figures 6.3.a-d show the WAXS patterns of sepiolite, PP and nanocomposites at

different filler concentrations. We can see that the sepiolite spectra has a prominent

peak at 2θ=7.2° which corresponds to the primary diffraction of the (100) crystalline

plane. The pattern of pure PP shows five main peaks, in the 2θ range of 10-30°,

140


characteristic of the monoclinic α-form [1]. For PP+Sep nanocomposites (Fig. 6.3.a)

we can observe that, except the reflection at 2θ=7.2°, enhanced for higher filler

content, which is typical of sepiolite, the same peaks of pure polymer are observed,

suggesting that mainly α-form is present. In particular there is no evidence of

reflections at 2θ=16°, which corresponds to the (300) plane of the β-phase, and at

2θ=20.3° corresponding to the characteristic (117) plane of the γ–phase.

5 10 15 20 25 30

a) PP+Sep

5%

2.5%

1%

PP

Sep

Inte

nsity

2 θ

5 10 15 20 25 30

b)

5%

2.5%

1%

PP

Sep

PP+PP-g-MA+Sep

Inte

nsity

2 θ

5 10 15 20 25 30

c) PP+PP-acid+Sep

5%

2.5%

1%

PP

Sep

Inte

nsity

2 θ

5 10 15 20 25 30

d) PP+Sep-sil

5%

2.5%

1%

PP

Sep

Inte

nsity

2 θ

Figure 6.3. X-ray diffraction spectra of: a) PP+Sep; b) PP+PP-g-MA+Sep; c)

PP+PP-acid+Sep; d) PP+Sep-sil nanocomposites at different concentrations of

filler, compared with virgin PP and pure sepiolite.

An interesting feature lies in the relative intensities of α-phase reflections. The

intensity of the peak at 2θ=17°, which corresponds to the (040) plane of α-phase,

141


increases with filler concentration while the peak at 2θ=14°, corresponding to the

(110) plane of α-phase, decreases. This can indicate a preferential orientation of PP

crystals induced by the nanofiller, with (040) planes parallel to the specimen surface

and the b-axes perpendicular to it [2-4]. Such phenomenon is ascribed to an epitaxial

growth of PP crystals from sepiolite surfaces. It is stressed here that such crystal

orientation relative to the sample is a direct consequence of a preferential orientation

of sepiolite nanoclay itself in the specimen. In the thin-film nanocomposite samples,

sepiolite is believed to be oriented in-plane with the film (see Paragraph 8.3.3). The

same conclusion may be drawn for the system PP+PP-g-MA+Sep in Fig. 6.3.b. A

different situation instead is observed with PP-acid or when Sep-sil is used instead

(Fig. 6.3c-d), with no evidence of a similar orientation enhanced by the nanofiller.

The DSC traces presented in Fig. 6.4 clearly show an increase of the crystallisation

temperature (Tc), implying that the sepiolite acts as a nucleating agent for the

crystallisation of polypropylene, but with significant differences for different

systems.

100 110 120 130 140

a)10%Sep

5%Sep

2.5%Sep

1%Sep

PP

Hea

t Flo

w [W

/g]

Temperature [°C]

100 110 120 130 140

b)10%Sep

5%Sep

2.5%Sep

1%Sep

PP

H

eat F

low

[W/g

]

Temperature [°C]

Figure 6.4. DSC traces corresponding to the non-isothermal crystallisation of a)

PP+PP-g-MA+Sep, b) PP+PP-acid+Sep at different filler content. The exothermic

peaks shift towards higher temperature as a result of the filler nucleating effect.

As can be also seen from Fig. 6.5, the systems with a poor distribution of the nano-

fibres (PP+Sep and PP+PP-g-MA+Sep) show the largest and continuous increase in

the crystallisation temperature while the systems with well dispersed sepiolite show

142


only a small effect and a limiting concentration of crystallisation nuclei is reached at

little sepiolite content after which no significant changes in Tc are observed.

From both WAXS and DSC experiments, it was observed so far in the paragraph that

there is a substantial difference in crystallisation activity between the systems that

present a good dispersion of sepiolite in PP and those characterised by agglomeration

of the filler.

0 2 4 6 8 10

110

115

120

125

130

135

PP+Sep PP+PP-g-MA+Sep PP+PP-acid+Sep PP+Sep-sil

T c ons

et [C

]

Filler [wt%]

Figure 6.5. Onset Temperatures of starting crystallisation in function of the filler

concentration for: ■ PP+Sep; ○ PP+PP-g-MA+Sep; PP+PP-acid+Sep;

PP+Sep-sil. A larger and continuous increase in the crystallisation temperature is

observed for PP+Sep and PP+PP-g-MA+Sep while a limiting concentration of

crystallisation nuclei is reached at 1 wt.% of filler for PP+PP-acid+Sep and

PP+Sep-sil.

According to some literature [5-7] any filler, including sepiolite, can nucleate PP on

its surface or at the connecting lines of two particles. In this case, the decrease of

nucleating efficiency coupled with a better dispersion of the particles seems to be a

143


strong indication for the second mechanism, i.e. that aggregates are the predominant

nucleating sites. But the sepiolite nucleation can be explained more simply by the

surface nature of the nanofiller and by the modification of its surface energy in

presence of a compatibiliser. Fillers have high-energy surfaces and can adsorb the

polymer preferentially along their crystal structure, acting as heterogeneous nuclei

for polymer crystals. When the filler surface is covered by an organic substance, the

surface free energy is decreased. This results, on one hand, in a reduction of

sepiolite-sepiolite inter-particle attraction and hence in a better nanoclay dispersion

but, on the other hand, on a reduced nucleating efficiency. This effect depends

specifically from the organic substance covering the inorganic surface. The two

functionalised polymers used in this chapter are then expected to be adsorbed on the

sepiolite in a peculiar way and influence differently the polymer nucleation. For

instance, PP-g-MA has comparably a low degree of functionalisation and the

molecular weight of the PP block in PP-g-MA is much higher than in the case of PP-

acid accounting for a weaker tendency to segregate to the filler surface. The

difference with Sep-sil is even more striking since in this case the alkyl-silane

coating is partially covalently bonded on the sepiolite surface.

Summarising, sepiolite nanoclay acts as an efficient nucleating agent for PP, which

affect both the onset of crystallisation temperature (PP starts crystallising at lower

temperatures) and the epitaxial crystal growth of PP on the sepiolite surface. The two

phenomena are closely connected and the last is the consequence of the former since

if there is heterogeneous nucleation of PP induced by the nanoclay this seems to

develop with an epitaxial growth of PP crystals on sepiolite surface. Moreover it was

demonstrated that such crystallisation activity can be altered by changing the surface

energy of the clay (i.e. by organo-modifying the clay surface), which has direct

impact on the sepiolite dispersability in PP.

144


6.3.3 Rheological Behaviour

The rheology of particulate suspensions in a polymer melt is governed by factors

such as the structure, size, shape and orientation distribution of the dispersed phase

but also the strength of interaction with the dispersant phase. Understanding the

rheological properties of polymer/clay nanocomposites is crucial in gaining insights

into the processability of these materials. Rheometry can be also a powerful tool for

investigating the nanocomposite morphology and how this relates to the final

properties. The rheological properties of PP/Sep-sil nanocomposites, with varying

loadings of the fibrous clay, are presented in this section (Fig. 6.6-8).

0.01 0.1 1 10 100

100

101

102

103

104

105

Frequency [Hz]

G' [

Pa]

100

101

102

103

104

105

G''

[Pa]

Figure 6.6. Frequency sweep test on polypropylene at 200°C.

At the temperature of 200 °C and the frequency range of 10-2-102 Hz, PP exhibits

(Fig. 6.6) the usual behaviour of homopolymer melts with low Mw distribution (i.e.

G’∝ω2 and G’’ ω). The crossover of the curves G’ and G’’ separates a liquid-like

behaviour at low frequencies from a solid-like behaviour at high frequencies. The

complex viscosity (Fig. 6.7) and the storage modulus (Fig. 6.8) increase, as expected,

with the addition of sepiolite clay.

∝

145


0.01 0.1 1 10 100

103

104

105

Com

plex

Vis

cosi

ty [P

a·s]

Frequency [Hz]

PP PP+1% Sep-sil PP+2.5% Sep-sil PP+5% Sep-sil PP+10% Sep-sil

Figure 6.7. Complex viscosity of PP/Sep-sil nanocomposites as a function of the clay

loading.

The increase relative to the virgin polymer, though, is much higher at low

frequencies, while it completely converges at high frequencies, for concentrations as

high as 2.5 wt.%. The zero-shear viscosity plateau (Fig. 6.7) shifts towards lower

frequency regions until it disappears, in the frequency range scanned, for filler

loadings of 10 wt.%. All the nanocomposites display significantly reduced frequency

dependence of the storage modulus, G’. The terminal slope of G’ rapidly decreases

as the clay loading increases and the nanocomposite filled with 10 wt.% of sepiolite

shows almost a plateau at low frequency range (Fig. 6.8), which is characteristic of

materials with solid-like behaviour (G’, G’’∝ω0). This can be explained if it is

assumed that the clay content has reached a threshold value to form a percolated clay

network [8], in which the sepiolite are incapable of freely rotating or moving. In this

case the mobility and relaxation of polymer chains would be seriously retarded in the

conned space created by the needle-like nanoclay. Such a percolation threshold is

then (Fig. 6.7-8) between 5 wt.% and 10 wt.% of filler. Although the calculation of

146


an accurate percolation threshold would require more experimental data (in particular

above the percolation) to be fitted with a suitable percolation theory, the range of

values given (5-10 wt.%) is compatible with the one predicted by Sheng et al. [8]. In

fact a weight fraction of 3-4 wt.% of attapulgite sticks, of aspect ratio 30-40, was

predicted to induce a percolating network. The slightly higher value observed for

sepiolite can be explained by the lower aspect ratio (average: s=27): in fact the

percolation threshold is generally inversely proportional to the aspect ratio.

0.01 0.1 1 10 100100

101

102

103

104

105

106

G' [

Pa]

Frequency [Hz]


Figure 6.8. Storage modulus of PP/Sep-sil nanocomposites as a function of the clay

loading.

6.3.4 Thermal Behaviour

The effect of different amounts of sepiolite on the thermal degradation of PP was

evaluated by TGA, performed both in inert atmosphere (N2) and in air. The results, in

accordance with recent publications [9-13], are presented in Fig. 6.9-10. In nitrogen

(Fig. 6.9), the degradation of PP is favoured at lower temperatures by the presence of

147


filler, except than at the very early stages of the heating process. The interpretation of

this phenomenon is commonly a catalytic effect of the sepiolite on the pyrolysis of

PP. The beneficial effect at the beginning of the degradation is probably due to

adsorption of PP volatile products by the microporous clay.

200 300 400 500 600

0

1

2

3

4

0

20

40

60

80

100

466 °C

479 °C

Der

ivat

ive

Mas

s [%

/°C]

Temperature [°C]

Mas

s [%

]


Figure 6.9. TGA of PP and PP+Sep-sil nanocomposites with different amounts of

nanoclays, in N2.

However, in oxidative atmosphere the thermal degradation of PP is substantially

retarded by the presence of sepiolite and the effect is more evident for increasing

loadings of the nano-filler, up to 5 wt.% (Fig. 6.10).

148


0

20

40

60

80

100

200 300 400 500 600

0.0

0.3

0.6

0.9

1.2

1.5

1.8

Mas

s [%

]


459 °C

399 °C

Der

ivat

ive

Mas

s [%

/°C]

Temperature [°C]

Figure 6.10. TGA of PP and PP+Sep-sil nanocomposites with different amounts of

nanoclays, in air.

The onset temperature of starting degradation is very similar among different

samples, while the temperature relative to the maximum degradation rate (peak of the

derivative of mass loss: Tmax) of PP/sepiolite nanocomposite, with 5 wt.% of filler, is

increased by 60 °C compared with the pure polymer. This can be explained, in

analogy with platelet-like nanoclays [14], with a barrier effect induced by the

sepiolite to the diffusion of oxygen. As the decomposition and volatilisation of PP

proceeds, the condensed phase is progressively enriched in inorganic phase,

149


preventing further diffusion of oxygen and thus retarding the degradation. The

previous explanation is strengthened by the fact that the Tmax in air of the PP/sepiolite

nanocomposite with 5 wt.% of filler is very close to the Tmax in inert atmosphere,

while these values differ by 80 °C for pure PP. In other words, the degradation

process of PP/sepiolite nanocomposites in oxidative atmosphere is shifted towards

degradation in inert atmosphere, because of the shielding effect induced by the

sepiolite.


The mechanical behaviour of the different nanocomposites is displayed in Fig. 6.11-

12 by representative stress-strain curves.

0.0 0.1 0.2 2 4 6 8 10

0

10

20

30

40

PP PP+1% Sep PP+PP-g-MA+1% Sep PP+PP-acid+1% Sep PP+1% Sep-sil

Stre

ss [M

Pa]

Strain

Figure 6.11. Stress-strain curves of different nanocomposites with 1 wt.% of sepiolite.

150


0.0 0.2 2 4 6 8 10

0

10

20

30

40

PP PP+5% Sep PP+PP-g-MA+5% Sep PP+PP-acid+5% Sep PP+5% Sep-sil

Stre

ss [M

Pa]

Strain

Figure 6.12. Stress-strain curves of different nanocomposites with 5 wt.% of sepiolite.

As expected, an increase in Young’s modulus is observed for all nanocomposite

systems investigated (Fig. 6.13).

0 1 2 3 4 5 61300

1400

1500

1600

1700

1800

1900

2000

2100

2200

2300


Youn

g's

Mod

ulus

[MPa

]

Filler wt. [%]

Figure 6.13. Young’s modulus of PP nanocomposites at different filler loadings.

151


Analogously, an enhancement in yield stress can be noticed in Fig. 6.14, with the

better performance for PP/Sep-sil nanocomposites.

0 1 2 3 4 5 627

28

29

30

31

32

33

34

35

36

37


Yiel

d St

ress

[MPa

]

Filler wt. [%]

Figure 6.14. Yield stress of PP nanocomposites at different filler loadings.

More interesting is the comparison of the strain at break for the different systems, as

shown in Fig. 6.15. While nanocomposites without pristine clay or with PP-g-MA

are dramatically embrittled at sepiolite concentrations as low as 2.5 wt.%, the use of

PP-acid and Sep-sil preserves the ductile nature of the polymer matrix, that undergo

yielding with necking stabilization and cold drawing even at filler concentrations

above 5 wt.%.

From the tensile tests presented in this paragraph, it appears that the nanocomposites

made by alkyl-silane functionalised sepiolite give the best mechanical performances,

in particular for what concern the yield stress.

152


0 1 2 3 4 5 6

0

200

400

600

800

1000 PP+Sep PP+PP-g-MA+Sep PP+PP-acid+Sep PP+Sep-sil

Stra

in a

t bre

ak [%

]

Filler wt. [%]

Figure 6.15. Strain at break of PP nanocomposites at different filler loadings. While

nanocomposites with pristine clay and with PP-g-MA undergo a clear embrittlement,

the use of PP-acid and Sep-sil preserves ductility even at filler concentrations above

5 wt.%.

6.3.6 Micromechanical Models

The results of the tensile tests, relative to PP/Sep-sil and PP/PP-acid/Sep

nanocomposites, shown in the paragraph before will be benchmarked with relevant

results from the scientific literature and interpreted in terms of the micromechanical

models of Halpin-Tsai and Pukanszky, for what concerns, respectively, the Young’s

modulus and yield stress.

Table 6.2 gives some comparative information about PP/clay nanocomposites

reported in the literature, among the most cited and influential in the field [15-17]. In

particular there are two sets of data relative to synthetic hectorite nanocomposites

modified with protonated primary hexadecyl amines. They differ in the use of PP-

MA as a compatibiliser with the PP matrix [15]. Other two sets of data from more

recent publications [16, 17] on PP/MMT nanocomposites are also included.

153


Table 6.2. Relevant PP/smectite clay nanocomposites reported in the literature.

Clay Compatibiliser Sample Preparation Ref

Type Treatment Type wt. [%]

ME1 C162 - - Injection Moulding [15]

ME C16 PP-MA3 20 Injection Moulding [15]

MMT DMDHTA5 PP-MA4 15 Compression Moulding [16]

MMT CPCl6 PP-MA7 20 Compression Moulding [17] 1 Synthetic hectorite (Somasif® ME100). 2 Protonated primary hexadecyl amine. 3 Hostaprime® HC5, from Hoechst AG (MA content 4.2 wt.%, Mn=4000, Mw/Mn=8). 4 From Homan Chemical Co (Korea) (Mw=59000, Mw/Mn=2.3). 5 Dimethyl dehydrogenated tallow ammonium anions. 6 N-cetyl pyridinium chloride. 7 Lycomont AR 504, from Clariant GmbH (MA content 3.5 wt.%, Mw=24000).

Young’s Modulus

The Halpin-Tsai equations, already introduced in Chapter 5, are used to interpret the

enhancement in stiffness of the PP nanocomposites with the clay content. The values

of Young’s moduli for the different set of data taken from the literature and from this

thesis are presented Figure 6.16.

It can be noticed that the Young’s modulus increases in a linear fashion, for filler

contents higher than a minimum value. The unfilled matrix, instead, often shows a

stiffness value which is incongruent with the rest of the data points. This can be

explained by an effect of polymer matrix modification (i.e. enhanced crystallinity or

change in crystal phase) induced by minute amounts of clays. In these cases, for

modelling purposes, the modulus of the matrix (E0) is extrapolated from a linear fit

of the data of composites with different filler content.

154


0 2 4 6 8 10

1500

2000

2500

3000

3500 PP+Sep-sil PP+PP-g-MA+Sep PP+C16-ME [15] PP+PP-g-MA+C16-ME [15] PP+PP-g-MA+MMT [16]

Yo

ung'

s M

odul

us [M

Pa]

Filler wt. [%]

Figure 6.16. Young’s modulus of PP/clay nanocomposites in function of the filler

wt.%. The lines are liner fits of the experimental data, excluding the pure matrix (0

wt.%).

In Figure 6.17, the increment of nanocomposites Young’s moduli relative to the

matrix moduli (E/E0) is presented, along with the model predictions for different clay

shape (plate or fibre) and aspect ratios. In order to calculate the filler vol.% from the

weight percent, the density of PP, sepiolite and montmorillonite (or hectorite) is

assumed, respectively, 0.9 gcm-3, 2.2 gcm-3 and 2.8 gcm-3. The shape factors relative

to fibres (sepiolite) and plates (montmorillonite and hectorite) are in accordance with

Paragraph 5.3.5, Chapter 5. Solid and dotted lines are, respectively, Halpin-Tsai

predictions for PP/Sep and PP/MMT (or PP/ME) nanocomposites, considering an

elastic modulus, of all types of clay, equals to 200 GPa (see 5.3.4). Moreover, the

specimens manufactured by injection moulding are supposed to have a unidirectional

155


orientation (1D) of the filler, whilst specimens manufactured by compression

moulding are supposed to have 2D randomly oriented filler (see Table 6.2).

0 1 2 3 4 5

1.2

1.6

2.0

2.4

2.8

3.2

3.6

4.0

4.4

∞→1D, s∞→2D, s

2D, s=27

2D, s=12

2D, s=4

1D, s=8

1D, s=21

PP+Sep2 PP+PP-acid+Sep PP+C16-ME [15] PP+PP-g-MA+C16-ME [15] PP+PP-g-MA+MMT [16]

E c/Em

Filler vol. [%]

Figure 6.17. Relative Young’s modulus versus filler vol.%. The lines are prediction

from the Halpin-Tsai equations for PP/Sep nanocomposites (fibre-like filler; solid

line) and PP/smectite clay nanocomposites (plate-like filler; dotted line), using true

filler aspect ratios (s) as a fitting parameter. In the graph the prediction for fibre-like

and plate-like filler for s ∞→ , corresponding to the rule of mixtures, are also

included. The abbreviations ‘1D’ and ‘2D’ stand for uniaxially oriented and 2D in-

plane randomly distributed filler, respectively.

The highest increase in elastic moduli is relative to the two PP/ME nanocomposites.

The use of PP-MA helps the platelet-like clays exfoliation, with an increased

reinforcement. However this result is to be explained mainly by the higher

orientation of the filler induced by injection moulding.

156


A way to compare the different sets of data is by fitting them with the Halpin-Tsai

equation and extrapolating the true aspect ratio (s) as a fitting parameter.

It can then be observed that the sepiolite based nanocomposites have an effective

fibre aspect ratio of s=12, which is lower than the average value estimated from TEM

micrograph in Paragraph 5.3.1 (s=27). This can be explained by a partial

agglomeration of sepiolite clay in PP (Fig. 6.1) and by a partial reduction in fibre

length induced by melt compounding (Fig. 6.2). Nevertheless, the true aspect ratio

for PP/Sep nanocomposites results to be higher than the aspect ratios of PP/MMT

nanocomposites (s=4) and PP/C16-ME nanocomposites (s=8) and close to the value

of s=21 relative to PP/PP-g-MA/C16-ME nanocomposites.

Since the aspect ratios expected from smectite clays are about 100-1000, it can be

appreciated how inefficiently plate-like clays are exfoliated into PP matrix and how

poor is their effective reinforcement. It can then be concluded that sepiolite is

comparably a more efficient nanofiller since the level of dispersion achievable is

higher.

Tensile Stress

An increase in stiffness alone is not a sufficient proof for reinforcement, since

Young’s modulus is expected to increases upon addition of inorganic fillers. Another

important mechanical property to take into account is the tensile yield stress that

depends on the components properties and particle dimensions and shape but also,

very importantly, on the filler/matrix interphase and the strength of interaction.

The results of tensile yield stress for the PP/Sep nanocomposites, from Figure 6.14,

can be interpreted in terms of the Pukanszky equation [18, 19] and compared with

the relevant scientific literature. The model assumes that an interphase forms

spontaneously in composites and yield stress changes proportionally to its actual

value as a function of composition. The composition dependence of tensile yield

stress can be described by the following equation:

157


)exp(5.21

10

ϕϕ

ϕσσ Byy +−

= Equation 6.1

where σy and σy0 are respectively the yield stress of the composite and the matrix, φ is

the volume fraction of the filler in the composites and B is a parameter related to the

load carried by the dispersed component and depends on interaction filler/matrix.

The term (1- φ)/(1+2.5 φ) expresses the effective load-bearing cross-sectional area of

the matrix. Without any interaction, the entire load is carried by the polymer and the

load-bearing cross-sectional area decreases with increasing filler content (the non-

interacting fillers act as holes). The value of the parameter B depends on all factors

influencing the load-bearing capacity of the filler, i.e. on the strength of interaction

and on the size of the contact surface. The effect of these factors on B is expressed as

0

ln)1(y

iyff lAB

σ

σρ+= Equation 6.2

where Af is the specific surface area of the filler (contact surface), ρf is its density,

while l and σyi are the thickness and the yield stress of the interphase. The latter two

parameters depend on the strength of the matrix/filler interaction.

If the model is valid, a linear correlation should be obtained when the natural

logarithm of reduced yield stress, σred (defined by Equation 6.3) is plotted against the

filler content (from Equation 6.1):

ϕϕσσ

−+

=1

5.21yred Equation 6.3

ϕσϕϕσσ Byyred +=⎟⎟⎠

⎞⎜⎜⎝

⎛−

+=

0ln

15.21lnln Equation 6.4

or, in terms of relative yield stress, σrel (defined by Equation 6.5):

158


ϕϕϕ

σσ

σ By

yrel =⎟

⎟⎠

⎞⎜⎜⎝

⎛

−+

=1

5.21lnln0

Equation 6.5

ϕϕϕ

σσ

σ By

yrel =⎟

⎟⎠

⎞⎜⎜⎝

⎛

−+

=1

5.21lnln0

Equation 6.6

Fig. 6.18 shows the tensile yield stress plotted with the filler volume fraction.

0 2 4 624

26

28

30

32

34

36

38

40

42

44

PP+Sep-sil PP+PP-acid+Sep PP+C16-ME [15] PP+PP-g-MA+C16-ME [15] PP+PP-g-MA+MMT [17]

Yiel

d St

ress

[MPa

]

Filler vol. [%]

Figure 6.18. Tensile yield stress of PP/clay nanocomposites as a function of the filler

vol.%.

Differently than the behaviour of Young’s modulus in Figure 6.16, the yield stress

doesn’t always show a monotone increase with the filler content. In analogy with the

previous paragraph, instead, the unfilled matrix often shows a stress value which is

incongruent with the rest of the data points and is excluded for modelling purposes.

This can be explained again by an effect of polymer matrix modification (i.e.

159


enhanced crystallinity or change in crystal phase) induced by minute amounts of

clays. The natural logarithm of the relative yield stress for PP/clay nanocomposites is

plotted against the filler volume fraction in Figure 6.19. From the linear fitting of the

experimental data points, a coefficient B can be extracted from the five sets of

experimental data points.

0 1 2 3 4 5

0.0

0.1

0.2

0.3

0.4

0.5

B=3.1

B=2.7

B=5.7

B=6.5

B=11.9

PP+Sep-sil PP+PP-acid+Sep PP+C16-ME [15] PP+PP-g-MA+C16-ME [15] PP+PP-g-MA+MMT [17]

Ln (σ

rel)

Filler vol. [%]

Figure 6.19. The natural logarithm of relative tensile stress of PP/clay

nanocomposites in function of the filler volume percent. From the linear fit of the

experimental data, the parameter B can be extracted.

The coefficient B varies between 2.7 and 11.9 (normally ranging from negative

values up to about 20 as a maximum [20]), which shows how sensitive this parameter

is in PP-based composites. For what PP/ME nanocomposites are concerned, for

instance, a parameter B increases four folds when PP-MA is added to the

nanocomposite. It is claimed here though, that the driving factor for such an increase

in B is more related to the difference in exfoliation/dispersion state of the nanoclays

160


rather than a significant change in the strength of the polymer/filler interface which

remains limited. In other words, referring to Equation 6.2, the change in B is related

more to Af (specific surface area of the filler) than to l and σyi. The statement above is

confirmed by the observation that PP+C16-ME nanocomposites are characterised by

a lower B parameter (Fig. 6.19) but a higher increase in Young’s modulus than

PP/Sep nanocomposites (Fig. 6.17). The higher stiffness is explained with the higher

orientation of the filler, induced by injection moulding, while the lower parameter B

is compatible with a poor exfoliation of the platelet clay and hence a lower aspect

ratio (from Fig. 6.17, s=8 is predicted) and effective contact surface.

6.4 Conclusions

Two different strategies have been followed in order to improve the dispersion and

compatibility of needle-like sepiolite nanoclay in isotactic propylene matrix: the use

of functionalised polymers as a third phase in the PP/Sep composites and the direct

functionalisation of the sepiolite surface. For what concern the former, a novel

copolymer (PP-acid) with high acid number (30) and lower molecular weight

(Mn=1800 gmol-1), was compared with the commonly used PP-g-MA, showing better

efficiency in terms of nanofiller dispersion and nanocomposites mechanical

properties. It is believed that an even better efficiency is achievable with further

optimizations of the thermodynamics of the surface-active polymers. The silanisation

of the sepiolite surface, on the other hand, is a very promising route to improve the

compatibility between PP matrix and the nanoclay. It guarantees excellent nanofiller

dispersion (up to 1-2 wt.%) and the best mechanical performances among the

different systems studied, in particular for what concern the yield stress.

In general, as for many other types of nanofillers, sepiolite considerably changes the

kinetics of crystallisation of PP, acting as heterogeneous nuclei and increasing the

crystallisation temperature. Moreover, although sepiolite did not alter the crystalline

structure and crystal polymorphism, it did seem to promote an orientation of the α-

phase crystals.

161


The addition of sepiolite into PP also significantly modify the rheology of PP melts,

reducing the frequency dependence of G’ and inducing a solid-like response for 10

wt.% of clays, which was explained by the formation a percolated network.

It was observed that the thermal degradation of PP, in oxidative atmosphere, is

substantially retarded by the presence of sepiolite and the effect is more evident for

increasing loadings of the nano-filler, up to 5 wt.%.

The mechanical performances of the best systems (PP+Sep-sil and PP+PP-acid+Sep)

are analysed with two micromechanical models (Halpin-Tsai and Pukanszky) and

compared with relevant publications in the scientific literature on more commonly

studied platelet-like clay nanocomposites. Sepiolite seams to provide a good

reinforcement in terms both of Young’s modulus and yield stress when compared

with smectite clays with equivalent level of filler orientation. The reason lays in the

better dispersion achievable with fibre-like rather than platelet-like clays instead of

specific interactions filler/matrix which are anyway very limited in the case of PP as

matrix.

6.5 References

1. G. Natta and P. Corradini, Structure and Properties of Isotactic Polypropylene.

Nuovo Cimento Suppl, 1960. 15: p. 40.

2. E. Ferrage, F. Martin, A. Boudet, S. Petit, G. Fourty, F. Jouffret, P. Micoud, P.

De Parseval, S. Salvi, C. Bourgerette, J. Ferret, Y. Saint-Gerard, S. Buratto, and

J.P. Fortune, Talc as nucleating agent of polypropylene: morphology induced

by lamellar particles addition and interface mineral-matrix modelization.

Journal of Materials Science, 2002. 37(8): p. 1561-1573.

3. F. Rybnikar, Orientation in Composite of Polypropylene and Talc. Journal of

Applied Polymer Science, 1989. 38(8): p. 1479-1490.

4. L.H. Wang and J. Sheng, Preparation and properties of polypropylene/org-

attapulgite nanocomposites. Polymer, 2005. 46(16): p. 6243-6249.

162


5. A. Pozsgay, T. Frater, L. Papp, I. Sajo, and B. Pukanszky, Nucleating effect of

montmorillonite nanoparticles in polypropylene. Journal of Macromolecular

Science-Physics, 2002. B41(4-6): p. 1249-1265.

6. B. Pukanszky and E. Fekete, Aggregation tendency of particulate fillers:

Determination and consequences. Polymers & Polymer Composites, 1998.

6(5): p. 313-322.

7. P. Svoboda, C.C. Zeng, H. Wang, L.J. Lee, and D.L. Tomasko, Morphology

and mechanical properties of polypropylene/organoclay nanocomposites.


8. L. Shen, Y.J. Lin, Q.G. Du, W. Zhong, and Y.L. Yang, Preparation and

rheology of polyamide-6/attapulgite nanocomposites and studies on their

percolated structure. Polymer, 2005. 46(15): p. 5758-5766.

9. J.L. Ahlrichs, C. Serna, and J.M. Serratosa, Structural Hydroxyls in Sepiolites.

Clays and Clay Minerals, 1975. 23(2): p. 119-124.

10. A. Marcilla, A. Gomez, S. Menargues, and R. Ruiz, Pyrolysis of polymers in

the presence of a commercial clay. Polymer Degradation and Stability, 2005.

88(3): p. 456-460.

11. G. Tartaglione, D. Tabuani, and G. Camino, Thermal and morphological

characterisation of organically modified sepiolite. Microporous and

Mesoporous Materials, 2008. 107(1-2): p. 161-168.

12. G. Tartaglione, D. Tabuani, G. Camino, and M. Moisio, PP and PBT

composites filled with sepiolite: Morphology and thermal behaviour.

Composites Science and Technology, 2008. 68(2): p. 451-460.

13. J.L. Valentin, M.A. Lopez-Manchado, P. Posadas, A. Rodriguez, A.M. Marcos-

Fernandez, and L. Ibarra, Characterization of the reactivity of a silica derived

from acid activation of sepiolite with silane by Si-29 and C-13 solid-state NMR.

Journal of Colloid and Interface Science, 2006. 298(2): p. 794-804.

14. R. Krishnamoorti and A.S. Silva, Polymer–clay nanocomposites, ed. T.J.

Pinnavaia and G.W. Brall. 2000, New York: Wiley.

15. P. Reichert, H. Nitz, S. Klinke, R. Brandsch, R. Thomann, and R. Mulhaupt,

Poly(propylene)/organoclay nanocomposite formation: Influence of

163


compatibilizer functionality and organoclay modification. Macromolecular

Materials and Engineering, 2000. 275(2): p. 8-17.

16. C.M. Koo, J.H. Kim, K.H. Wang, and I.J. Chung, Melt-extensional properties

and orientation behaviors of polypropylene-layered silicate nanocomposites.

Journal of Polymer Science Part B-Polymer Physics, 2005. 43(2): p. 158-167.

17. L. Szazdi, B. Pukanszky, G.J. Vancso, and B. Pukanszky, Quantitative

estimation of the reinforcing effect of layered silicates in PP nanocomposites.

Polymer, 2006. 47(13): p. 4638-4648.

18. B. Pukánszky, B. Turcsányi, and F. Tüdős, Effect of Interfacial Interaction on

the Tensile Yield Stress of Polymer Composites, H. Ishida, Editor. 1988,

Elsevier: New York. p. 467-477.

19. B. Turcsanyi, B. Pukanszky, and F. Tudos, Composition Dependence of Tensile

Yield Stress in Filled Polymers. Journal of Materials Science Letters, 1988.

7(2): p. 160-162.

20. L. Szazdi, A. Pozsgay, and B. Pukanszky, Factors and processes influencing

the reinforcing effect of layered silicates in polymer nanocomposites. European

Polymer Journal, 2007. 43(2): p. 345-359.

164

7 Oriented PP / Sepiolite Nanocomposite Tapes

7.1 Introduction

This chapter presents the study of solid-state drawn nanocomposites tapes based on

iPP and sepiolite clay, a field in which the literature is not much investigated [1-4].

In this case, the intrinsic anisotropy of sepiolite nano-fibres, allows also an

exploitation of the orientation of such filler upon drawing. The nanocomposites tapes

are fully characterised for their morphology, structure, orientation and mechanical

properties. In particular, extraordinary enhancements in mechanical properties are

found to be provided by small amounts sepiolite (<2.5 wt.%).

7.2 Experimental

7.2.1 Materials

The iPP used was a homopolymer polypropylene resin H507-03Z from DOW (MFI

3.2 g/10min at 230 °C/2.16 Kg). The sepiolite Pangel® were supplied by Tolsa

165

CHAPTER 7 – Oriented PP / Sepiolite Nanocomposite Tapes

(Spain) and surface-modified with alkyl-silane at TNO (The Netherlands). For

simplicity the filler will be referred to as Sep throughout the chapter, although it was

differently addressed to in the previous one, in order to be distinguished from the

pristine clays.

7.2.2 Composite Tape Preparation

Rectangular specimens (50x3 mm2) were cut from thin films (100 µm in thickness),

prepared by compression moulding (220 °C for 10 min) granules of pre-compounded

nanocomposite at different concentrations (see Chapter 6). The desired draw ratio

was obtained by solid state drawing the initial rectangular specimen (Fig. 7.2) at a

temperature of 120 °C in an universal testing machine (Instron 5584) equipped with

an environmental chamber. The temperature of 120 °C is intermediate between the

melting point (~165°C) and the crystalline relaxation temperature Tα (~100°C) of

polypropylene [5, 6].

λ=1

λ=2

λ=5

λ=24

3mm50mm

Figure 7.1. Schematic illustration of nanocomposite tape preparation. A rectangular

specimen is cut from a 100 µm thick compression moulded film and drawn in the

solid state to a tape of the desired draw ratio.

166


7.2.3 Composite Tape Characterisation


Morphological studies were carried out using Scanning Electron Microscopy (SEM)

nalysis (Jeol JSM-6300F) on gold-coated surfaces. The lateral surfaces of oriented

investigated after ductile failure induced by tensile tests.

CD camera, in transmission mode, with an exposure time of 30sec. The X-ray

as a synchrotron radiation at the EPRF (European Synchrotron Radiation

nanocomposite tapes samples at room

mperature in an Instron 5586, equipped with a 1kN load cell and Merlin data

tware. A crosshead displacement rate of 5 mm/min was used on

a

nanocomposite tapes were

WAXS

Two-dimensional wide angle X-ray diffraction (WAXS) patterns were recorded by a

C

source w

Facility), beam line BM26 Dubble (Dutch-Belgian beam line). The average

wavelength was 1.24 Å. The two-dimensional X-ray patterns were transformed into

one-dimensional patterns by performing integration of the Azimuthal intensity. All

data manipulation on WAXS results were performed with the data analysis program

FIT2D, available free for academic users

Tensile Tests

Tensile tests were conducted on the

te

acquisition sof

rectangular samples with a length of 50 mm. The thickness and width, depending on

the draw ratio, were in the range of 100-30 µm and 0.8-3 mm, respectively. The

width was measured from optical images obtained using an Olympus BX60F and

recoded with a JVC colour camera, model KY-F55BE. The thickness was measured

167


with a digital micrometer with resolutions of 1 µm. For all the samples at least five

specimens were measured and the mean value was calculated.

DSC

Differential scanning calorimetry (DSC) was performed on about 1 mg samples of

anocomposite tapes in standard aluminium pans, using a Mettler-Toledo DSC 822e.

temperature up to 220 °C at a scanning rate of 10

7.3.1 Morphology of Tapes

igure 7.2 shows the SEM micrographs of the planar surface of PP tapes and

sepiolite. The tapes, of an original draw ratio

f 20, were imaged after ductile failure induced by tensile tests. PP tapes are

n

Sample were heated from room

K/min, held at that temperature for five minutes and then cooled at 10 K/min to room

temperature. The degree of crystallisation of the nanocomposite tapes was calculated

from the integration of the endothermic peak in the first heating scan.


F

composites with 1 wt.% and 10 wt.% of

o

characterised by a highly oriented and micro-fibrillated structure (Fig. 7.2.a-b). The

same considerations can apply for the PP+1%Sep composite tapes (Fig. 7.2.c-d), for

which the addition of 1 wt.% of filler doesn’t substantially change the overall picture.

Instead, a different morphology is shown from the PP+10%Sep composite tapes (Fig.

7.2.e-f), in particular for what concerns a reduced micro-fibrillation (Fig. 7.2.e) and

168


the presence of inorganic inclusions. Therefore the incorporation of as much as 10

wt.% inorganic filler negatively affects the orientation of the tapes.

a)

b)

c)

d)

e)

f)

Figure 7.2. SEM micrographs, at different magnifications, of the lateral surface of:

a-b) PP tapes, c-d) PP+1%Sep tapes and e-f) PP+10%Sep, after ductile failure. A

white circle indicates sepiolite agglomerations.

169


Moreover, as reported in the previous chapter (Fig. 6.1.d-e), sepiolite loadings above

5 wt.% are scarcely well dispersed in polypropylene matrix and will partially

agglomerate at a micron level. These agglomerates can act as defects and stress

oncentration points which seriously compromise the stability of the tapes upon

anical properties of polymers are closely related to the polymer molecular

Many experimental techniques have been employed to determine the

scale. In this study wide angle X-

y diffraction (WAXD) has been used to investigate the polymer crystals

Figure 7.3. Scheme of WAXS measurements in the through direction.

c

drawing.

7.3.2 PP and Sepiolite Orientation

The mech

orientation.

degree of orientation in polymers on a molecular

ra

orientation. WAXS measurements have been taken in the through (out-of-plane)

direction of various nanocomposite tapes, as schematically shown in Figure 7.3.

X-ray

Drawing direction

170


2D wide angle X-ray images of PP tapes and PP/sepiolite nanocomposite tapes at

different draw ratios are presented in Figure 7.4. Each of the patterns is individually

optimised in contrast and brightness. The so-called Debye-Scherrer rings are

observed for film samples (λ = 1), which is a clear indication of random orientation.

Each ring represents a crystal plane of polymer crystal cell. Five main reflections are

observed in Fig. 7.4.a, which are typical of the PP α crystal form [7] as already

discussed in the previous chapter (Fig. 6.3). In Fig.7.4.c-d an extra ring appears

which is characteristic of the basal crystal plane of the sepiolite. Upon drawing, the

crystallites will align along the drawing direction and this can be associated to the

change in the X-ray pattern, in which the rings become arcs. At a draw ratio of 20

(Fig. 7.4.b-d-f), crystallites are perfectly aligned and the diffraction arcs are very

narrow, almost dots. It is important to underline at this stage that a full understanding

of the kind of preferred orientation that prevails in a selected sample could not

simply be obtained from a WAXS test with the fibre axis perpendicular to the direct

beam. The intensities of appropriate diffractions should be observed as a function of

the angle of tilt of the fibre axis. Furthermore, when the mode of preferred

orientation is more complex than simple axial, the intensities of certain reflections

are found to depend not only on the inclination but also on the azimuth of the

specimen. In the case of non-orthogonal crystal systems the study is even more

complicated by the fact that there are no crystalline planes normal to the chosen

crystallographic direction (usually the c-axis: the direction of polymer

macromolecules). In the case of i-PP, the α-form crystal cell is monoclinic and it can

be demonstrated that two independent sets of planes need to be evaluated to fully

describe the orientation of the polymer chains with respect to one direction [8].

171


λ=1

λ=20

a)

b)

d)c)

e)

f)

Figure 7.4. 2D WAXS patter

tapes, at λ=1 (left column)

indicate the main reflection p

(131)(041)

(111)
(110)
n

a

la

(

(040)

s

nd

n

10

(130)

of: a-b) PP, c-d) PP+5%Sep and e-f) PP+10%Sep

λ=20 (right column), respectively. Black arrows

es of PP and of sepiolite clays.

0)clay

172


However, there is an easier approach to this, which is widely used. In fact, there is a

correspondence between the intensity in a diffraction arc at a particular angle to the

meridian and the number of crystallites oriented at the same critical angle to the fibre

axis, as was firstly pointed out by Weissenberg [9]. When the fibre axis is

perpendicular to the incident beam it can be shown that, for moderate to small Bragg

angles (θ~0), the angle of inclination of the plane normal to the fibre axis (φ ) is

practically equal to the Azimuthal angle β as directly measured from the meridian

. If a me e of a quate intensity can be obtained, the

degree of the molecular chain, taken to be coincident with the c-axis, is given

directly by the Azimuthal extension of that plane. However, since meridional

reflections are usually weak, near-meridional crystalline planes can be used or even

equatorial crystalline planes as an indirect index of degree of alignment.

In this study, as often the case in the polymer field, an indicative measurement of

crystallite orientation with respect to the drawing direction, will be obtained from

sity as a function of

on the film ridional (00l) plan de

Azimuthal scans of the (110) diffraction plane, which appears as the most intense.

If the WAXS 2D images (Fig.7.4) are integrated along the 2θ axis, plots as in

Figure7.5.a-b are obtained, showing the integrated X-ray inten

the Azimuthal angle φ for PP and PP+5%Sep tapes with draw ratios of 1, 7 and 20

respectively. The background intensity is subtracted for all samples and the 0 o

represents the equatorial direction. It is observed that un-stretched films show

isotropic behaviour of polymer chains which could be determined from the constant

diffraction intensity over the whole angle range in all films. When films are stretched

under solid-state condition, the intensity increases and the peak becomes sharper,

which indicates that the drawing leads to h gher degree of orientation of the polymer

chains in both samples.

i

173


-45 -30 -15 0 15 30 45

a) PP, λ=1 PP, λ=7 PP, λ=20

Inte

nsity

[a.u

.]

Azimuthal Angle [degree]

b)

-45 -30 -15 0 15 30 45

Inte

nsity

[a.u

.]

Azimuthal Angle [degree]

PP+5% Sep, λ=1 PP+5% Sep, λ=7 PP+5% Sep, λ=10

Figure 7.5. X-ray intensity (integrated along the 2θ axis) versus the Azimuth angle

for a) PP sample and b) PP+5%Sep tapes at λ=1, 7 and 20. Solid line represents

Gaussian fitting of the data points.

174


Assuming a uniaxial Gaussian distribution, the Azimuthal peaks are fitted with a

Gaussian curve and the widely accepted Hermans’ orientation factor [10] could be

calculated for both the drawn PP and PP/sepiolite tapes. The crystal orientation

function fc is calculated from the Hermans’ orientation equation, which takes the

form:

( ) φφφ22 sin

2311cos3

21

−=−=f Equation 7.1

Where φ2cos and φ2sin represent respectively the mean-square cosine and sine

(averaged over all the crystallites) of the angle between a given crystal axis and the

fibre axis, which serves as the reference direction.

For a given crystallographic axis, φ2cos assumes values of 1 for perfect

alignment, 1/3 for random orientations, and 0 for precise perpendicularity. At the

same time assumes the respective values 1, 0 and -1/2.

In case of axial orientation with respect to a given direction Z (i.e. the drawing

direction), numerical values of the mean-square cosines, of the angles that the (hkl)

plane normals form with Z, can be calculated from the intensity distribution I (φ):

φf

∫

∫= 2/

0

2/

0

2

,2

sin)(

cossin)(cos π

π

φφφ

φφφφϕ

dI

dI

Zhkl Equation 7.2

175


where ф is the Azimuthal angle, at 90o corresponds to the drawing direction, and

ф=0 o corresponds to the transverse direction of the tapes. Figure 7.6 shows the

calculated Hermans’ orientation factors obtained from WAXD experiments for

composite tapes at different draw ratios. It is interesting to note that Hermans’ factor

of PP polymer crystals at draw ratio of 7 is already approaching 1, the maximum in

e Hermans’ factor range of values.

0.8

1.0

th

1 3 5 7 9 11 13 15 17 19 21

n Fa

ctor

0.6

0.0

0.2

0.4

PP PP+1% Sep PP+2.5% Sep PP+5% Sep PP+10% Sep

nto

igure 7.6. Orientation of PP crystals: Hermans’ orientation factor as a function of

draw ratio λ, relative to PP tapes with different concentrations of sepiolite.

The result shows that solid-state drawing is an efficient method in obtaining

significant polymer alignment at relative low draw ratio. A second important

bservation lays in the similar values of the orientation factor for tapes of equivalent

draw ratios and different amount of filler. This affirms that the polymeric matrix in

Orie

ati

Draw Ratio λ

F

o

176


the nanocomposite tapes and in the pure PP tapes possess the same degree of

crystalline orientation at a given macroscopic draw ratio. This is important since it

allows an easier understanding and comparison of data to extract the effect of

different filler content on the mechanical properties of tapes of similar draw ratio.

The solid state drawing of the composite tapes not only orients the polymer crystals

but also leads to an alignment of the sepiolite needles.

1 3 5 7 9 11 13 15 17 19 21

0.0

0.2

0.4

0.6

0.8

1.0

Orie

ntat

ion

Fact

or

Draw Ratio λ

Figure 7.7. Orientation of sepiolite: Hermans’ orientation factor for PP+5%Sep

composite tape, as a function of draw ratio λ.

In analogy with the study of the polymer crystal orientation, an Azimuthal scan can

be performed on the basal diffraction ring (100) of the clay, at 2θ about 7.2 °, in

order to study the orientation of the needle-like clays.

177


Figure 7.7 shows the Hermans’ orientation factors for composite tape with 5 wt.% of

sepiolite at different draw ratios. It can be concluded that at a draw ratio of about 7

both the polymer crystals and the needle-like nanofillers are perfectly aligned along

the tape direction.


In this paragraph the mechanical properties of the oriented tapes are characterised by

tensile tests. Figure 7.9 shows the stress-strain curves of PP tapes at different draw

ratios.

100

150

200

250

300

350

400

450

500

550

600

0 5 10 15 20 250

50

400 600 800 1000

λ=24

λ=20

λ=16

λ=9

λ=1

Stre

ss [M

Pa]

%]

he mechanical performance of

polypropylene. Upon drawing to λ=24, the Young’s modulus increases from about

Strain [

Figure 7.8. Stress-strain curves of PP tapes of different draw ratios λ.

The solid-state drawing process dramatically changes t

178


1.3 GPa to 14 GPa and the ultimate tensile stress from about 50 MPa to 550 MPa. In

contrast, the strain at break is reduced to 5 %. These values are still far away from

the theoretical limit of PP, for which a stiffness of about 40 GPa is expected [11].

The reason is well established and lies in the fact that the drawability of melt-

rystallised PP achieved by solid-state drawing is limited by the presence of c

molecular entanglements [12-15].

16

18

4 6 8 10 12 14 16 18 20 22 24 26

4

6

8

10

12

14

s=27∞→s

s=12

Youn

g's

Mod

ulus

[GPa

]

Draw ratio λ

PP PP+2.5% Sep PP+10% Sep

Figure 7.9. Young’s modulus of nanocomposites tapes at different draw ratios. The

dotted lines are Halpin-Tsai predictions of PP tapes filled with 2.5 wt.% sepiolite,

completely aligned in the direction of the tape, at three aspect ratios: s=12, which

was fou aspect

ratio of sepiolite nanofibres (5.3.1), and s

nd to fit the isotropic samples (Fig. 6.15), s=27, which is the average

, which corresponds to the upper ∞→

bound of the rule of mixtures. The Halpin-Tsai equation is applied to the unfilled PP

tapes at different draw ratios upward in order to shift the elastic moduli values.

179


The effect of nanoclays on the mechanical performance of the tapes is presented in

Fig. 7.9-7.13. The addition of small amounts of sepiolite to PP (up to 2.5 wt.%)

significantly enhances the stiffness of the tapes for draw ratios as high as 20 (Fig.

7.9). For comparison, the Halpin-Tsai predictions of PP tapes filled with 2.5 wt.% of

sepiolite, fully oriented in the direction of the tapes, is also included in the same

graph (dotted lines). Three aspect ratios are considered: s=12, which is the fitted

value found for isotropic specimens in the previous chapter (see Fig. 6.15), s=27,

which is the average aspect ratio of sepiolite nanofibre (see Paragraph 5.3.1), and

, which corresponds to the rule of mixtures and represents the upper bound.

The Halpin-Tsai equation is applied individually to the unfilled PP tapes at each

draw ratio in order to shift the relative elastic modulus value upward. It is surprising

to observe that the experimental data for PP+2.5%Sep tapes, at intermediate draw

ratios (10<λ<20), lie above the upper bound predictions based on the

micromechanical model. PP+2.5%Sep tapes with extremely low or high draw ratios

(λ<7 or λ>24) instead, present moduli very close to the ones of pure PP tapes.

Therefore, in this last case sepiolite has negligible reinforcing effect on the polymer

matrix. On the other hand, nanocomposites tapes with loadings above 5 wt.% have

stiffness even inferior to those of PP tapes for the entire draw ratio range. This

ere t

s ∞→

variation with the filler content can be better represented in Fig. 7.10. In this graph

the Young’s moduli of tapes of the same draw ratio (λ=9, 16 and 20) are replotted as

a function of increasing sepiolite wt.%, showing a maximum value for systems based

on 2.5 wt.% of clay (Fig. 7.10). An analogous comment can be given for the strength

of the tapes shown in Fig. 7.11 and 7.12, wh he ultimate tensile stress reaches an

optimum value for composite tapes with a filler content of 2.5 wt.% (until λ=20).

180


0 2 4 6 8 104

6

8

10

12

14

16

Youn

g's

Mod

ulus

[Pa

]G

Filler wt [%]

λ=9 λ=16 λ=20

Figure 7.10. Young’s modulus of nanocomposite tapes of λ=9, 16 and 20, as a

function of sepiolite filler content.

6 8 10 12 14 16 18 20 22 24

200

250

300

350

400

450

500

550

UTS

[MPa

]

Draw Ratio λ

PP PP+2.5% Sep PP+10% Sep

Figure 7.11. Ultimate tensile stress of nanocomposites tapes in function of λ.

181


0 2 4 6 8 10

200

250

300

350

400

450

500

550

UTS

(MPa

)

Filler wt%

λ=9 λ=12 λ=20

Figure 7.12. Ultimate tensile stress of nanocomposites tapes in function filler wt.%.

6 8 10 12 14 16 18 20 22 24 265

6

7

8

9

10

11

12

13

14

15

Stra

in a

t bre

ak [%

]

Draw Ratio λ

PP PP+1% Sep PP+2.5% Sep PP+5% Sep PP+10% Sep

Figure 7.13. Strain at break of nanocomposites tapes.

182


There are no significant differences, instead, in the strain at break (Fig. 7.13). Failure

strain is dominated by the polymer with little effect of the sepiolite. Hence, it can be

summarised so far that PP tapes are exceptionally reinforced by the addition of small

amounts of nanoclays (best values at 2.5 wt.% of sepiolite) at intermediate draw

ratios (10<λ<20), without compromising the final strain at break. The reason for this

behaviour can not be explained by classical composite theories and further

investigations are required (see also Paragraph 7.3.4).

Nevertheless, the detrimental effects of high concentrations of sepiolite and high

draw ratios (λ>24) can already be further commented. For what concerns the first

aspect it can be explained in terms of two main factors: filler dispersion and physical

,

can not be easily dispersed in polypropylene. In this case

the clays will tend to agglomerate in micrometric bundles. Agglomerates in contrast

to individually dispersed sepiolite nanoclays, act as defects and stress concentration

points that prevent drawability and polymer alignment. With respect to the second

factor, it is helpful to refer to the rheological data in chapter 6. There it was shown

that the addition of clays increases the viscosity of PP/sepiolite melts but, for

sepiolite concentrations above 5 wt.%, there was a change in the viscoelastic

behaviour at low frequencies. A solid-like behaviour appears rather than a liquid-

like, which was attributed to a percolated physical network formed by the filler. This

means that at higher concentrations of sepiolite, the physical entanglements increase

so much that the drawability of polymer chain is seriously compromised.

Passing to analyse the negative effect of high draw ratio, instead, a primary

explanation is the imperfect

entanglements. As shown by the SEM micrographs in Chapter 6 and also in Fig. 7.2

high loadings of sepiolite

the debonding which can develop upon drawing due to

183


interfacial adhesion PP/clay (Paragraph 6.3.6). Another explanation relies on simple

nd extent it as a more general

hermal/fire

composite theories. At the beginning of the stretching process (up to λ=6-7), there is

the initial contribution from the orientation of the filler (Fig.7.7) on the matrix

reinforcement, which is hindered in this study since the draw ratios investigated are

above this threshold. However, upon further drawing, the reinforcing effect of

sepiolite progressively diminishes, since the mechanical properties of the PP matrix

considerably increases.

It is worth spending few more words on this point a

consideration. It was already mentioned that the theoretical stiffness of PP crystals is

about 40 GPa in comparison with 230 GPa of polyethylene [11]. The modulus of

sepiolite, as measured via nano-bending tests in chapter 5, is around 200 GPa. It is

obvious that it would not make any sense to try to reinforce ultradrawn PE fibres

with nanoclays since the Young’s modulus of the reinforcement is close to the values

of the matrix itself. Clearly, this is different for the case of PP. Well dispersed

sepiolite, if perfectly bonded to the polymeric matrix, is potentially an interesting

nanofiller for oriented PP tapes and fibres. It is also important to keep in mind that

the use of PP tapes or fibres is not only due to its mechanical properties. Other fibres,

such as PE, have intrinsically much higher performances. However, one of the most

interesting properties of PP, over PE, is its comparatively high melting temperature.

With the addition of small amounts of nanoclay is possible to even further increase

the utilisation temperature of PP tapes. In fact rheological results have demonstrated

that the flowability properties of PP, at a certain temperature, are significantly

reduced. Another interesting result of adding sepiolite is the improved t

resistance. In Chapter 6 it has been demonstrated, via thermo gravimetric analysis,

how PP/sepiolite composites have a significant delay in the degradation temperature.

184


7.3.4 Thermal Analysis

The increase in orientation upon solid-state drawing can be indirectly followed by

the increase in crystallinity of the different tapes. The degree of crystallinity is

mH∆ ) obtained from the first heating scan of the

:

calculated from the melting heat (

DSC measurements, according to the following equation

cm

c fH

X ×H∆

∆=

0

Equation 7.3

where 0H∆ =209 J/g is the melting enthalpy of 100 % crystalline PP and cf is a

factor that simply takes in account for the excluded mass of inorganic that does not

melt. For instance, cf is 0.95 for a 5 wt.% nanocomposite. For this specific

investigation the second heating scan of DSC does not give much information. After

melting and re-crystallisation, the oriented structure of the tape is lost and an almost

isotropic sample is obtained, as already characterised in Chapter 6. The results are

plotted in Figure 7.15, for different composite tapes as a function of draw ratio. The

nanocomposite tapes, for intermediate values of λ, have a higher crystallinity

compared with the pure PP tapes. As the draw ratio increases, the crystallinity

decreases on a relative scale and this behaviour is more pronounced for higher

concentrations of sepiolite (>5 wt.%). Figure 7.14 shows remarkable similarities with

the mechanical results presented in Figures 7.9.

185


45

55

60

65

70

2 4 6 8 10 12 14 16 18 20 22 2425

30

35

40

50

75

PP PP+1% Sep PP+2.5% Sep

X c]

PP+5% Sep

[%

Draw io λ

Figure 7.14. Degree of crystallinity of different nanocomposite tapes in function of λ.

Rat

If the Young’s modulus is plotted in function of the crystallinity of the PP tapes (Fig.

7.15), a good correlation is found between the two parameters and the sets of data

points relative to composite tapes with different filler loadings, seam to lie on the

same line. This observation suggests that the deviations of the properties of

nanocomposite tapes from the pure PP tapes (Fig.7.9 and 7.11), are partially due to

changes in polymer crystallinity and semi-crystalline structure induced by the

presence of nanoclay during solid-state drawing, rather than a direct reinforcement of

the sepiolite.

186


30 35 40 45 50 55 60 65 70 750

2

4

6

8

10

12

14

16

Youn

g's

Mod

ulus

[GPa

]

Xc [%]

PP PP+1% Sep PP+2.5% Sep PP+5% Sep

Figure 7.15. Young’s modulus of nanocomposites tapes in function of the degree of

AXS studies. Small amounts of sepiolite (< 2.5 wt.%) have an

extraordinary effect on the mechanical properties of PP tapes, for intermediate draw

ratios (10<λ<20). The DSC analysis suggests that the origin of the improved

mechanical properties is more related to the modification of the semi-crystalline

polymer structure rather than a pure reinforcement effect. Nevertheless, sepiolite

polymer crystallinity.

7.4 Conclusions

In this chapter, the effect of sepiolite nanoclay on the final properties of oriented

polypropylene tapes has been evaluated. Solid-state drawing effectively orients both

the polymer crystallites and the needle-like filler at relatively low draw ratios, as

shown by W

187


remains an interesting nanofiller for oriented polypropylene tapes. Beside improved

mechanical behaviour, sepiolite clay provides increased utilisation temperature

(lower flowability) and improved thermal and fire resistance.

7.5 References

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Science, 2004. 92(1): p. 552-558.

2. M. Joshi and V. Viswanathan, High-performance filaments from

compatibilized polypropylene/clay nanocomposites. Journal of Applied

A. Gerasin, B.F. Shklyaruk, L.A.

Tsamalashvili, H.R. Fischer, and I.V. Razumovskaya, Structure and

r of nanocomposites based on polypropylene and modified

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Polymer Science, 2006. 102(3): p. 2164-2174.

3. E.M. Antipov, A.A. Barannikov, V.

deformation behavio

4. E.M. Antipov, M.A. Guseva, V.A. Gerasin, Y.M. Korolev, A.V. Rebrov, H.R.

Fischer, and I.V. Razumovskaya, Structure and deformation behavior of

nanocomposites based on LDPE and modified clays. Polymer Science Series

A, 2003. 45(11): p. 1130-1139.

5. B. Alcock, N.O. Cabrera, N.-M. Barkoula, C.T. Reynolds, L.E. Govaert, and

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188


6. T. Schimanski, T. Peijs, P.J. Lemstra, and J. Loos, Influence of Postdrawing

Temperature on Mechanical Properties of Melt-Spun Isotactic Polypropylene.


7. G. Natta and P. Corradini, Structure and Properties of Isotactic Polypropylene.

uppl, 1960. 15: p. 40.

Alexander, X-ray diffraction methods in polymer science. 1969, New

9.

10.

11. -ray diffraction method to the

12. prehensive Composites, T.-W.

13. I.M.

Nuovo Cimento S

8. L.E.

York.: Wiley.

K. Weissenberg, Z. Physik, 1961. 8: p. 20.

P.H. Hermans and P. Platzek, Kolloid Z., 1939. 88: p. 68.

K. Nakamae and T. Nishino, The application of X

measurement of crystal deformation and crystal modulus of high polymers.

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T. Peijs, M.J.N. Jacobs, and P.J. Lemstra, Com

Chou, A. Kelly, and C. Zweben, Editors. 2000, Elsevier Science Publishers

Ltd: Oxford. p. 263-302.

Ward, Structure and properties of oriented polymers. 1997, London:

Chapman & Hall.

14. P. Smith and P.J. Lemstra, Ultrahigh-Strength Polyethylene Filaments by

Solution Spinning-Drawing .2. Influence of Solvent on the Drawability.

Makromolekulare Chemie-Macromolecular Chemistry and Physics, 1979.

180(12): p. 2983-2986.

15. P. Smith, P.J. Lemstra, and J.P.L. Pijpers, Tensile-Strength of Highly Oriented

Polyethylene .2. Effect of Molecular-Weight Distribution. Journal of Polymer


189

8 Polyamide 6 / Sepiolite

Nanocomposites

8.1 Introduction

The first successful report of polymer-clay nanocomposites was, as mentioned

earlier, obtained via in situ intercalative polymerisation of ε-caprolactam, in the

presence of platelet-like clays [1, 2]. Although there are already commercial

applications of this new material, the simple melt compounding is much more viable

and attractive from a commercial viewpoint. In this chapter the preparation of

PA6/sepiolite clay nanocomposite, as obtained in a mini twin-screw extruder, will be

described, as well as a characterisation of the material in terms of morphology (via

SEM and TEM analysis), crystallinity and crystal structure of the semi-crystalline

matrix and mechanical performances. The tensile tests results will be interpreted in

terms of two micromechanical models, the Halpin-Tsai model for stiffness and the

Pukanszky equation for strength, and compared with the results on PA6/MMT of

selected publications.

190

CHAPTER 8 – Polyamide 6 / Sepiolite Nanocomposites

8.2 Experimental

8.2.1 Materials

Polyamide 6 (PA6) Technyl® C206 from Rhodia (density of 1.14 gcm-3) and

sepiolite Pangel® from Tolsa (Spain) were used in this study.

8.2.2 Nanocomposites Preparation

Nanocomposites were prepared by a two-step blending process in a mini twin-screw

extruder DSM Micro 15, at 240 °C and 200 rpm for 10 min, utilising a continuous

flow of nitrogen gas into the mixing chamber, to minimise possible polymer

degradation. First, sepiolite was mixed with the PA6 to make a masterbatch at 20

wt.% of filler, which was subsequently diluted with neat PA6 to obtain nanofiller

concentrations of 0.1 wt.%, 1 wt.%, 2.5 wt.% and 5 wt.%. Both polymer pellets and

sepiolite powder were dried overnight at 80°C before being compounded.

Tensile test specimens were prepared by compression-moulding nanocomposite

pellets, previously compounded in mini-extruder, at 240 °C for 5 min under a

constant force of 40 kN, followed by a cooling step to room temperature. The

compression-moulding was realised with a Benchtop Press Rondol, and the tensile

test specimens mould was in-house made according to standard dimensions (ASTM

D-638).



Morphological studies were carried out using a Jeol JSM-6300F Scanning Electron

Microscope (SEM), on gold-coated, cold-fractured samples. The surface analysed

was the cross-section of the tensile test specimens and the brittle fracture was

191


obtained after immersion in liquid nitrogen. Transmission Electron Microscopy

(TEM) analyses were realised with a Jeol JEM 2010 microscope, on ultra-thin

samples obtained by a microtome.

WAXS

X-ray diffraction photographs were taken using the Cu Kα radiation, generated with

a RINT-2000 (Rigaku Co.), at 40 kV and 20 mA. The distance between sample and

detector was 37.5 mm.

DSC


Calorimeter (DSC) Mettler-Toledo 822e. All samples (~4 mg) were firstly heated to

250 °C and kept at that temperature for 5min to remove any thermal history and then

cooled at a rate of 10 K/min to analyse non-isothermal crystallisation behaviour. A

successive scan at 10 K/min until 250 °C was utilised to study the melting of the

composite samples and the overall crystallinity, given that the heat of fusion for the

completely crystalline PA6 is 240 J/g.

TGA

Thermo Gravimetric Analyses were performed with a TA instrument Q500 on about

10mg samples cut from tensile test specimens. The test was performed at a scanning

rate of 20 K/min up to 900 °C, under inert atmosphere (N2) conditions.

Tensile Tests





standard.

192




As already stated in previous chapters, a good dispersion of the nanofiller into the

polymeric matrix is precondition to an efficient reinforcement.

a)

b)

c)

d)

e)

f)

Figure 8.1. SEM micrographs of: a)-b) PP+1%Sep, c)-d) PP+2.5%Sep, e)-f)

PP+5%Sep, at magnification of 5000 and 10000 times respectively. A good

dispersion of sepiolite nanoclays is evident even at relative high filler content.

193


In this paragraph a series of SEM and TEM micrographs are presented in order to

discuss the morphology of PA6/sepiolite nanocomposites. Figure 8.1.a-f refers to

cold fractured surfaces of nanocomposites samples with different filler

concentrations (1 % to 5 %). Two magnifications are presented for each

concentration (5000 and 10000 times). We can observe an excellent distribution of

inorganic filler in its finest elemental units even at concentration as high as 5 % in

weight.

Figure 8.2, instead, shows two TEM micrographs of the PA6+5%Sep

nanocomposite. Even at the high magnifications obtained by Transmission Electron

Microscopy, no indication of nano-filler aggregation is present.

a) b)

Figure 8.2. TEM micrographs of PA6/sepiolite nanocomposites with 5 wt.% of filler

at different magnifications.

We can conclude that polyamide 6 is a promising polymeric matrix for dispersing

sepiolite nanoclays even at relative high concentrations of filler and without the need

to employ specific surface functionalisations or compatibilisers, which was the case

194


for polypropylene (Chapter 6). The polarity of PA6 can account for its compatibility

with hydrophilic nanoclays, such as sepiolite.

8.3.2 Non-Isothermal Crystallisation

In this section the influence of sepiolite concentration on the non-isothermal

crystallisation behaviour of PA6 is studied. Figure 8.3 presents the DSC traces

relative to the first cooling scan of PA6 and his nanocomposites.

140 160 180 200 220

Exo

10% Sep

5% Sep

2.5% Sep

1% Sep

PA6

Heat

Flo

w [W

g-1]

Temperature [°C]

Figure 8.3. DSC crystallisation peaks of PA6/sepiolite nanocomposites at different

filler concentrations.

All the nanocomposites samples have higher crystallization temperatures (Tc) than

the pure PA6 sample, which can be attributed to a nucleating effect induced by the

nanoclay [3-5]. With regards to the effect of clay concentration, it is noticed that the

195


crystallisation temperature of the polymer slightly decreases with increasing amount

of the filler. The explanation of this behaviour is that the introduction of more

particles hinders chain mobility and retards crystal growth [3].

Figure 8.4 shows the effect of addition of sepiolite on the melting of polyamide

matrix as taken from the second heating scan.

180 200 220 240

Endo

10% Sep

2.5% Sep

5% Sep

1% Sep

PA6

Heat

Flo

w [W

g-1]

Temperature [°C]

Figure 8.4. DSC melting peaks of PA6/sepiolite nanocomposites at different filler

concentrations.

For neat PA6, only one endothermic peak is observed at 221 °C. A second peak at

lower temperature (215 °C) appears, as a shoulder to the first peak, when clays are

added to the matrix. The formation of this second peak has been associated, by

several authors, to a less stable crystalline phase (γ phase) of PA6 [6-8]. The lower

melting point of the γ-phase is believed to be due to the lower crystalline density and

increased entropy of melting compared to the α-phase [9]. Although this hypothesis

196


is realistic, the appearance of low temperature melting peaks as well as broadening

effects has also been ascribed to changes in crystallite thickness and its distribution

[3]. Integrating the endothermic peaks of Fig. 8.4 and taking the heat of fusion for

the completely crystalline PA6 as 240 J/g [3], the sample crystallinity was calculated

and reported in Figure 8.5.

0 2 4 6 8 1025

26

27

28

29

30

31

32

33

34

35

X c [%

]

Filler wt [%]

Figure 8.5. Amount of crystalline phase for PA6/sepiolite nanocomposites. The heat

of fusion for the completely crystalline PA6 is taken as 240 J/g [3]. Virgin PA6 is

presented, for comparison, as open circle, while full squares refers to processed

samples.

The addition of clay results in a minor decrease in crystallinity, when compared with

the processed PA6. The presence of high concentrations of sepiolite prevents the

formation of large crystalline regions due to spatial hindrance and to specific

interaction between polymer and filler. This leads to smaller crystallite structures and

197


less perfect crystalline lamella. Such imperfections may also explain the lower

melting points observed for the nanocomposites (Fig.8.4).

8.3.3 WAXS - PA6 Crystal Structure

Nanocomposite samples were also analysed with the WAXS technique in order to

validate hypotheses made by the study of the DSC thermograms in the previous

paragraph. 2D wide-angle X-ray measurements have been taken in the through

direction, shown schematically in Figure 8.6, on tensile test dog-bone specimens.

X-ray

Figure 8.6. Scheme of WAXS measurements in the through direction.

The results of the measurements are shown in Figure 8.7. For neat PA6 (Fig.8.7.a),

two main diffraction rings were observed at 2θ ~ 20 ° and 23.7 ° which are

commonly associated, respectively, with the (200) and (002/202) crystal planes of

the α-phase of polyamide 6 [9-11]. The α-phase population was the dominant

crystalline phase in virgin PA6.

At relative smaller angles (2θ ~ 7.2 °), the typical basal diffraction of sepiolite clays

was observed for PA6/sepiolite nanocomposite samples (see also Chapter 5 and 7),

with the intensity increasing with clay concentration (Fig.8.7.b-c).

Two new diffraction rings also appeared, with addition of sepiolite, at 2θ ~ 10.7 °

and 21.3 ° (Fig.8.7.b-c). The latter two were attributed to the (020) and (002) crystal

planes of the γ-phase, respectively. Therefore, the addition of clay induced the

formation of a different crystal form of polyamide 6 (γ-phase) as was also suggested

from the DSC results (Fig.8.4).

198


a) b)

c)

(200)α (002/202)α

(002)γ

(020)γ

(100)clay

Figure 8.7. WAXS through view: a) PA6, b) PA6+1%Sep, c) PA6+5%Sep. White

arrows show the principal diffraction rings with the corresponding crystal planes.

The stabilisation of γ crystal phase in PA6 has already been reported by early

investigations on platelet-like clays [6, 12]. This phenomenon is particularly

important since a change in crystal structure may affect physical and mechanical

properties of the polymeric matrix. Ito et al. [13] showed that the stiffness of the α-

phase is higher than the γ-phase below the glass transition temperature, but it

decreases more rapidly at higher temperatures. This results in a higher heat distortion

temperature of the γ-phase. The same authors [13] also reported that the draw ratio

199


achieved for PA6 fibres drawn at temperature between 110 °C and 180 °C is higher

if the crystal is in the γ-form.

It is worthy mentioning that the complex crystallinity and polymorphism of PA6/clay

nanocomposites may not necessary be induced only by the presence of clays. Zapata-

Espinosa et al. [14] found that the molecular weight was the main cause of the

thermal behaviour of in situ polymerised PA6/clay hybrids, with a minimum

influence of clay concentration and molecular orientation. Moreover it is well known

that the crystal structure obtained when PA6 is crystallised from the melt is

influenced, among others, by thermal conditions, applied stress, presence of moisture

and additives. For instance, rapid cooling and low crystallisation temperature

promote the γ-form of PA6, while higher crystallisation temperature and slow

cooling leads to the α-form [15-18].

WAXS measurements have also been taken in the edge direction, as schematically

shown in Figure 8.8, on tensile test dog-bone specimens and are presented in Figure

8.9.

X-ray

Figure 8.8. Scheme of WAXS measurements in the edge direction.

Similar conclusions can be drawn, from the edge view WAXS experiments

concerning the crystal structure. The addition of sepiolite induces the formation of γ-

form of PA6, as evident by the appearance of the two diffraction rings at 2θ ~ 10.7 °

and 21.3 ° (Fig. 8.9.b-c).

200


a) b)

c)

(100)clay

(100)clay

Figure 8.9. WAXS, edge view: a) PA6, b) PA6+1%Sep and c) PA6+5%Sep. White

arrows shows orientation of sepiolite, as can be seen from the diffraction at 2θ ~ 7.2 °.

It is interesting to observe the basal diffraction of the sepiolite, at 2θ ~ 7.2 °. The full

ring, still present in the through view (Fig. 8.7.b-c), was transformed in two arc-like

diffractions along the equatorial direction. This phenomenon can be explained by

partial orientation of sepiolite nano-clays in the longitudinal plane of the specimen.

In order to better visualise the results obtained from WAXS investigation, a 3D

rendering of a tensile test specimen is provided in Figure 8.10.

201


a)

b)

Figure 8.10. Schematic 3D image of a nanocomposite tensile test specimen where

the nanofiller is aligned in-plane. Sepiolite nano-fibres are represented in white and

are not in scale with the specimen dimensions. The real length of sepiolite clay is

about 4000 times smaller than in the picture.

The top view (Fig. 8.10.a) represents a completely random orientation of the fibre-

like filler, while the edge view (Fig. 8.10.b) shows how sepiolite clays are

preferentially oriented in-plane with the sample. This particular micro-structure is the

result of the sample preparation method employed that, as already mentioned,

involves compression moulding of pre-compounded nanocomposite pellets.

202


8.3.4 Thermo Gravimetric Analysis

The thermo gravimetric analyses are presented in Figure 8.11. The pure PA6 sample

and the nanocomposites behave in a very similar fashion in inert atmosphere. The

presence of sepiolite induces a minor delay in the polymer thermal degradation as it

is seen from the small shift of the curves towards higher temperature.

300 400 500 600

0

20

40

60

80

100 PA6 1% Sep 2.5% Sep 5% Sep

Wei

ght L

oss

[%]

Temperature [°C]

Figure 8.11. TGA of PA6/Sep nanocomposites, in inert atmosphere (N2).

A more significant effect is in the final residue at the end of the test, which increases

with the amount of clay. In this way, the actual content of filler can be accurately

obtained, once the thermal degradation of sepiolite itself is taken in account (refer to

Chapter 5). The results are summarised in Table 8.1.

203


Table 8.1. Filler content of nanocomposites samples as obtained from TGA.

Nominal filler wt. [%] TGA residue (at 800°C) [%]

PA6+0.1%Sep 0.1 0.40

PA6+1%Sep 1.0 1.45

PA6+2.5%Sep 2.5 2.58

PA6+5%Sep 5.0 5.23


The tensile mechanical performances of PA6 nanocomposites are displayed in Figure

8.12 by representative stress-strain curves.

0 10 20 30 40 500

10

20

30

40

50

60

70

80

90

100

110

120

Stre

ss [M

Pa]

Strain [%]

PA6 PA6+0.1% Sep PA6+1% Sep PA6+2.5% Sep PA6+5% Sep

Figure 8.12. Stress-strain curves of PA6/Sep nanocomposites at different filler

loadings.

204


Figure 8.13 and 8.14 show the reinforcement effect of different sepiolite loadings in

polyamide 6, in terms of enhancement of Young’s modulus and tensile stress.

0 1 2 3 4 5 62.5

3.0

3.5

4.0

4.5

5.0

5.5

Youn

g's

Mol

udus

[GPa

]

Filler wt. [%]

Figure 8.13. Young’s modulus of PA6/sepiolite nanocomposites as a function of the

filler loading.

0 1 2 3 4 5 675

80

85

90

95

100

105

110

UTS

[MPa

]

Filler wt. [%]

Figure 8.14. Ultimate tensile stress of PA6/sepiolite nanocomposites as a function of

the filler loading.

205


Both stiffness and strength increase proportionally with filler content up to a

concentration of 5 wt.%. In particular, the stiffness of PA6 showed an increase of 60

% with about 5 wt.% of sepiolite clays. In order to have a term of comparison, a

similar level of reinforcement is obtained with as much as double the content in

weight of glass fibres, according to data sheets from Rhodia on commercial

reinforced PA6 [19]. Moreover the specimens were probably prepared by the

common injection moulding technique, which induces much higher level of filler

orientation and overestimate the filler reinforcement efficiency.

These results are particularly interesting, if compared with the same tests on

PP/sepiolite nanocomposites (Chapter 6). In polypropylene matrix, the reinforcement

induced by the filler was decreasing with increasing sepiolite concentration. For

loadings above 2.5 wt.%, even a reduction of stiffness was shown for some

composite formulations (Fig. 6.13). Instead, in the case of polyamide 6, the

reinforcement efficiency of the filler is high even for relatively high concentrations

of filler. This phenomenon can have two explanations. First, a good level of

dispersion of individual nano-clays is easily achieved in PA6 at concentrations as

high as 5 wt.%, as demonstrated by the SEM and TEM micrographs in Figure 8.1

and 8.2. It is stressed that this result has been obtained by a simple melt-

compounding technique, without any compatibiliser as was necessary for

polypropylene. The second reason is the good compatibility of the polar polyamide 6

and a hydrophilic filler such as sepiolite. Strong hydrogen bonding interactions are

expected to take place between the amide groups of PA6 and the characteristic

hydroxyl groups of the sepiolite [20]. In fact, without an effective interaction

between matrix and reinforcing phase, leading to the formation of an extended

interphase region, the filler mainly acts as an inclusion, decreasing the effective load-

bearing cross-sectional area of the matrix and, as such, reducing the matrix strength.

Finally, the strain at break of different nanocomposites is presented in Figure 8.15.

As often in the case of particle filled thermoplastic, the strain at break decreases with

increase of reinforcing phase. This reduction though is not as dramatic as in the case

of less compatible matrices such as polypropylene. At loading of 5 % in weight of

206


sepiolite, PA6 still preserves 1/7th of the initial strain at break, while un-

compatibilised PP completely looses its ductility (Fig 6.9).

0 1 2 3 4 5 60

10

20

30

40

50

St

rain

at b

reak

[%]

Filler wt. [%]

Figure 8.15. Strain at break for PA6/sepiolite nanocomposites in function of the

filler loading.

8.3.6 Micromechanical models

In analogy with previous chapters, the results of the tensile tests will be benchmarked

with relevant results from the scientific literature and interpreted in terms of the

micromechanical models of Halpin-Tsai and Pukanszky, for what concerns,

respectively, the Young’s modulus and ultimate tensile stress.

Table 8.2 lists same information relative to the clays and the sample preparations

employed for PA6 nanocomposites reported in scientific publications [21-24], among

the most cited and influential in the field.

207


Table 8.2. Relevant PA6/clay nanocomposites reported in the literature.

Clay Sample Preparation Ref

Type Treatment

1 MMT (USA)1 Ammonium Salt2 Injection Moulding [21]

2 MMT (JP)3 Ammonium Salt2 Injection Moulding [22]

3 MMT1 Ammonium Salt4 Injection Moulding [23]

4 MMT3 - ? [24]

5 Sep - Compression Moulding 5

1 Clay source: Wyoming, USA. Commercial name: Cloisite Na+, by Southern Clay Products. 2 Bis(2-hydroxy-ethyl)methyl tallow ammonium chloride. 3 Clay source: Yamagata, Japan. Commercial name: Kunipia-P, by Kunimine Industries. 4 Methyl, hydrogenated tallow, 2-ethylhexyl ammonium methylsulfate. 5 Data from this thesis.

In particular the data relative to PA6/Sep from this chapter are compared with the

best results (entries 1-2). These refer to melt compounded PA6/MMT

nanocomposites, obtained using organo-modified montmorillonite clays from two

well known mines: Yamagata (Japan) and Wyoming (USA) [21, 22]. The difference

between the two clays is that the former has slightly larger average particle length

and hence aspect ratio.

Young’s Modulus

The Halpin-Tsai equations, already introduced in Chapter 5, are used to interpret the

enhancement in stiffness with the clay content. The increment of nanocomposites

Young’s modulus relative to the matrix modulus is presented Figure 8.16, along with

the model predictions for different clay shape, aspect ratios and degree of orientation.

In order to calculate the filler volume fraction from the weight percent, the density of

PA6, sepiolite and montmorillonite were assumed respectively, 1.14 gcm-3, 2.2 gcm-3

and 2.8 gcm-3. The shape factors relative to fibres (sepiolite) and plates

(montmorillonite) are again in accordance with Paragraph 5.3.5, Chapter 5.

208


0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5

1.0

1.2

1.4

1.6

1.8

2.0

2.2

2.4

2.6

1D, s=45

∞→2D, s

∞→1D, s

1D, s=36

2D, s=50

2D, s=27

PA6/Sep PA6/MMT (USA) [21] PA6/MMT (JP) [22]

E c/Em

Filler vol. [%]

Figure 8.16. Relative Young’s modulus versus filler vol.%. The lines are prediction

from the Halpin-Tsai equations for 2D randomly oriented PA6/Sep nanocomposites

(fibre-like filler; solid line) and uniaxially (1D) oriented PA6/MMT nanocomposites

(plate-like filler; dotted line), using true filler aspect ratios (s) as a fitting parameter.

The condition corresponds to the upper bound predictions of the rule of

mixtures.

∞→s

The PA6/Sep nanocomposites data points lay below the ones relatives to the

PA6/MMT. In first analysis, this could be explained by the higher aspect ratios of

MMT clay. However it must be stressed that the PA6/MMT nanocomposites

specimen are obtained by injection moulding, which induces a higher degree of

orientation of the filler compared to compression moulding, as for PA6/Sep. Solid

lines are the Halpin-Tsai predictions relative to the sepiolite nanocomposites and are

drawn considering an elastic modulus of the clay of 200 GPa (see Paragraph 5.3.4),

and assuming a 2D random distribution of filler. The above hypothesis is supported

209


by the WAXS results (see Figure 8.9), which showed a partial in-plane orientation of

sepiolite fibre-like clays. It can be seen that the average aspect ratio of 27 slightly

underestimates the experimental data. A best fit is found for an aspect ratio of around

50. It is very encouraging to notice that the theoretical predictions for fillers of

infinite aspect ratio, corresponding to the rule of mixtures, is not much higher,

indicating that sepiolite in PA6 has a very high reinforcing efficiency.

The dotted lines, instead, are relative to MMT nanocomposites. The platelet-like clay

is assumed to have the same elastic modulus as the sepiolite. In accordance with Ref.

[22], the composite is considered uniaxial (because of the orientation induced by

injection moulding). The fitted aspect ratios are 36 and 45 (for the USA and Japan

mine), which are lower than the values extrapolated from TEM micrographs [22] (57

and 69, respectively) and, more importantly, lower than the values expected from

fully exfoliated MMT clays (100-1000). It is very surprising that the sepiolite aspect

ratio is even higher than the MMT ones. This tells us that MMT are not as efficiently

reinforcing the matrix as predicted from Halpin-Tsai equations.

Tensile Stress

The results of tensile strengths for the PA6/Sep nanocomposites, from Figure 8.14,

can be interpreted in terms of the Pukanszky equation [25, 26], which has already

been introduced in paragraph 6.3.6. In Figure 8.17, tensile stress is plotted with the

nanoclay vol.%. PA6/Sep nanocomposites show a linear correlation. The same is true

for PA6/MMT nanocomposites with the exception of the initial value (σ0) which is

excluded from modelling purposes. This can be explained by the polymer matrix

modification (i.e. enhanced crystallinity or change in crystal phases), induced by the

presence of MMT, which obviously can not be contemplated by the model.

210


0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.565

70

75

80

85

90

95

100

105

110

PA6/Sep PA6/MMT (USA) [21] PA6/MMT (JP) [22]

U

TS [M

Pa]

Filler vol. [%]

Figure 8.17. Ultimate tensile stress of PA6/clay nanocomposites in function of the

filler vol.%.

In Figure 8.18, the natural logarithm of the relative tensile stress for PA6/Sep and

PA6/MMT nanocomposites is plotted against the filler vol.%. From the linear fitting

of the experimental data points, a coefficient B can be extracted from the three sets of

experimental data points. It is reminded here that B is a parameter related to the load

carried by the dispersed phase (filler) and depends on interaction filler/matrix (see

Pag. 154). The coefficient B varies between 11 and 13, which can be considered as a

very good result if compared with values obtained for other systems, ranging from

negative values up to about 20 as a maximum [27]. This shows a remarkably good

interaction between PA6 and clay. Moreover, the parameter B relative to PA6/Sep, is

slightly higher that that of the two PA6/MMT nanocomposites. This result is

particularly interesting if compared to the Young’s modulus analysis, for which

211


PA6/Sep nanocomposites showed a smaller increase in stiffness with filler content

compared with PA6/MMT (which was explained by a different filler orientation).

0.000 0.005 0.010 0.015 0.020 0.025 0.030

0.00

0.05

0.10

0.15

0.20

0.25

0.30

0.35

0.40

PA6/Sep Melt PA6/MMT(USA) Melt [21] PA6/MMT(JP) Melt [22]

B=12.6

B=11.6

B=13.1

Ln (σ

rel)

Filler Vol. fraction

Figure 8.18. The natural logarithm of relative tensile stress of PA6/Sep and

PA6/MMT nanocomposites in function of the filler vol.%. From the linear fit of the

experimental data, the parameter B can be extracted.

Referring to equation 8.8 and given that the surface area, for exfoliated platelet-like

clay (Paragraph 5.3.2), and the density of MMT is higher than those of sepiolite, the

higher value of the coefficient B must be the consequence of a thicker and stronger

interphase ( l and iσ ) between sepiolite and PA6 rather than montmorillonite and

PA6. This can be explained by the silanol groups present on the sepiolite surface (but

not on MMT), which can strongly interact with the characteristic amide group of

PA6, via hydrogen bonds. For this reason also, no functionalisation was necessary

212


for sepiolite while montmorillonite were cation exchanged with quaternary

ammonium salts.

8.4 Conclusions

PA6/sepiolite nanocomposites have been successfully prepared by melt

compounding in a mini-extruder apparatus. The morphology, as investigated by SEM

and TEM, showed a homogeneous dispersion of the nanoclays into the polar

polymeric matrix. According to non-isothermal crystallisation tests conducted with

DSC, sepiolite didn’t nucleate the crystallisation of PA6 and didn’t increase the total

amount of crystallinity (slightly diminished). WAXS confirmed the appearance of γ-

phase crystals, against the more stable α-phase, induced by the presence of sepiolite

in PA6. Edge view X-ray tests also established partial orientation of sepiolite nano-

fibres in the longitudinal plane of tensile test specimen.

Notable enhancements in stiffness and strength were obtained with only little weight

percentage of filler although this was accompanied by a decrease in strain at break.

The mechanical performances were interpreted in terms of two micromechanical

models (Halpin-Tsai and Pukanszky equation) and compared with the results of

PA6/MMT taken from selected literature publications. The better reinforcement in

Young’s modulus induced by the platelet-like clay was attributed to the higher

degree of particle orientation.

Sepiolite, instead, showed a remarkable enhancement in tensile stress, explained by

strong hydrogen bonds between silanols groups on sepiolite surface and the amide

groups of the matrix, which are instead missing in montmorillonite even after surface

functionalisation.

213


8.5 References

1. A. Okada, M. Kawasumi, A. Usuki, Y. Kojima, T. Kurauchi, and O.





Research, 1993. 8(5): p. 1179-1184.



4. Y.P. Khanna, R. Kumar, and A.C. Reimschuessel, Memory Effects in

Polymers .3. Processing History Vs Crystallization Rate of Nylon-6

Comments on the Origin of Memory Effect. Polymer Engineering and

Science, 1988. 28(24): p. 1607-1611.

5. Y.P. Khanna, A barometer of crystallization rates of polymeric materials.

Polymer Enginnering and Science, 1990. 30(24): p. 1615-1619.

6. L.M. Liu, Z.N. Qi, and X.G. Zhu, Studies on nylon 6 clay nanocomposites by

melt-intercalation process. Journal of Applied Polymer Science, 1999. 71(7):

p. 1133-1138.

7. K. Varlot, E. Reynaud, M.H. Kloppfer, G. Vigier, and J. Varlet, Clay-

reinforced polyamide: Preferential orientation of the montmorillonite sheets

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216

9 In Situ Polymerised Polyamide 6 /

Sepiolite Nanocomposites

9.1 Introduction

After the successful preparation and characterisation of PA6/sepiolite

nanocomposites obtained by simple and industrial-friendly melt compounding

method (Chapter 8), this chapter will present two different nanocomposites formed

by in situ polymerisation [1, 2] of PA6 polymerised either in presence of unreactive

pristine sepiolite or grafted from amino-silane functionalised sepiolite. These

represent two distinct model cases, differing in the interfacial interactions between

matrix/filler. It will be demonstrated how the interphase is crucial for the final

properties of the nanocomposite.

9.2 Experimental

9.2.1 Materials

Two different sepiolite nanoclays were employed: a natural sepiolite (Pangel®) and

an aminopropyltriethoxysilane (Dynasylan HS2909 from Degussa) functionalised

sepiolite, both supplied by Tolsa (Spain). They will respectively be referred to as Sep

217

CHAPTER 9 – In-Situ Polymerised PA 6 / Sepiolite Nanocomposites

and Sep-NH2, from now on in this chapter. The silanisation of the nanoclay was

carried out by Tolsa following a reported procedure [3] and the amount of surface

functionalisation was quantified in 3 wt.%, from which it can be calculated a

concentration of 0.26 nmol of primary amine groups per gram of clay. Independent

researchers have disclosed that the silanes bonding covalently bonded through

condensation reaction with the silanol groups on the surface of sepiolite, were about

30 % while the rest was simply adsorbed [4, 5].

For the polymerisation reaction, the monomer (ε-caprolactam 99 %) and the initiator

(1-hexadecylamine) were purchased from Aldrich, while the catalyst (Phosphinic

Acid H3PO2) was purchased from Fluka. They were all used as received. The

commercial PA6 used for the compounding was the Technyl® C206 from Rhodia.

9.2.2 Masterbatches Preparation - Polymerisation

Two masterbatches, with 20 wt.% inorganic phase, were produced by in situ

polymerisation of ε-caprolactam in presence of pristine and amino functionalized

sepiolite. A typical procedure involved drying about 4 g of Sep-NH2 at 100 °C

overnight, followed by drying under vacuum at room temperature for 24 hours to

avoid the presence of water that could hydrolyse the monomer. 16g of ε-caprolactam

were added to the clays together with 0.037 ml of the catalyst H3PO2 (Mw= 66,

d=1.21 g/ml), in a ratio of 0.25 mol% relative to the monomer. The mixture, in a

sealed glass vessel and under vacuum, was initially heated up to 100 °C, in an oil

bath, for 1 hour and vigorously stirred to obtain a homogeneous suspension. The

polymerisation was carried out at 250 °C for 24 hours. In the case, above, of the

amino functionalised sepiolite the polymerisation is initiated directly from the

primary NH2 groups on the surface of the clay and thus can be described as a grafted-

from polymerisation.

For pristine sepiolite, the polymerisation was carried out following a similar method

as above. However, as unfunctionalised sepiolite doesn’t possess any chemical

groups able to initiate the polymerisation of ε-caprolactam, an external primary

amine, 1-hexadecylamine (Mw= 241), was introduced. The amount of initiator was

218


0.25 g for 16 g of monomer, which is equivalent to the moles of clay-bound amine in

the Sep-NH2 case.

All solid samples obtained from the above polymerisations were subsequently

pulverised at cryogenic temperature (in liquid N2) and unreacted caprolactam and

higher oligomers were removed by subsequent washing with hot distilled water (90

°C) for 24 hours.

The in situ polymerised PA6/Sep-NH2 masterbatch was subjected to washing with 99

% Formic Acid in order to roughly estimate the amount of grafted PA6 chains on the

sepiolite. 3 g of the composite were added to 80 ml of solvent followed by

centrifugation of the suspension to eliminate the supernatant. The procedure

described above was repeated three times, after which the weight percent of

inorganic was evaluated by TGA.

9.2.3 Masterbatches Dilution - Melt Compounding

The two masterbatches, previously prepared by in situ polymerisation, were diluted

with neat commercial PA6 to obtain nanofiller concentrations of 0.1 wt.%, 1 wt.%,

2.5 wt.%, 3 wt.% and 5 wt.%. The different nanocomposites were prepared by means

of a mini twin-screw extruder DSM Micro 15, operating at 200 rpm of screw speed

for 10 min at a temperature of 240 °C. Nitrogen gas was purged in the mixing

chamber during compounding in order to minimise the polymer degradation.

Polymer pellets and the in situ polymerised masterbatches powders were dried

overnight at 80 °C before compounding.


Gel Permeation Chromatography (GPC)

GPC sample preparation and tests were performed by RAPRA on the commercial

PA6, the in situ masterbatches, and the nanocomposites prepared after dilution of the

masterbatches. The solvent used was 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP). The

219


test specimen consisted of 5-10 mL of solution with a polymer concentration of

approximately 2 mg/mL. Solutions were left overnight and, after thorough mixing,

filtered through a 0.45 µm PTFE filter. A refractive index detector was used and the

data were collected and analysed with the Polymer Laboratories ‘Cirrus’ software.

The GPC set-up was calibrated with polymethylmethacrylate (PMMA) and the

results are expressed as ‘PMMA equivalent’ molecular weight. Samples are run in

duplicates.

TGA

Thermo Gravimetric Analyses were performed with a TA Instrument Q500 on

samples of about 10 mg. The test was performed at a scanning rate of 20 K/min up to

1000 °C, in inert (N2) or oxidative (air) atmosphere.


Morphological studies were carried out using a Jeol JSM-6300F Scanning Electron

Microscope, on gold-coated, cold-fractured samples. The surface analysed was the

cross-section of the tensile test specimens and a brittle fracture was induced after

immersion, for sufficient time, in liquid nitrogen. Transmission Electron Microscopy

(TEM) analyses were also performed with a Jeol JEM 2010 instrument, on ultra-thin

samples obtained by ultra-microtome.

DSC


Calorimeter (DSC), Mettler-Toledo 822e. All samples (~4 mg) were firstly heated to

250 °C and kept at that temperature for 5 min, to remove any previous thermal

history, and then cooled at a rate of 10 K/min to analyse non-isothermal

crystallisation behaviour. A successive scan at 10 K/min until 250 °C was utilised to

study the melting of the composite samples and the overall crystallinity, given that

the heat of fusion for the completely crystalline PA6 is 240 J/g.

220


Tensile Tests





standard. At least 10 specimens have been tested for each sample and the average

values and standard deviations calculated. Specimens were obtained by compression

moulding at 240 °C with a Rondol bench press and were dried at 80 °C under

vacuum overnight before being tested. The toughness of the specimens was

estimated by the integration of the area of the force-displacement tensile tests curves.

DMA

Dynamic Mechanical Analyses were carried out with a TA Instrument DMA Q800,

in tensile mode, equipped with film/tape clamps. The specimen were nanocomposite

tapes of dimensions 40x3x0.1 mm3, cut from thin films obtained by compression

moulding at 240 °C with a Rondol bench press. The tests were performed at a

heating rate of 3 K/min and constant frequency of 1 Hz. The specimens were dried at

80 °C under vacuum overnight before being tested.


9.3.1 Gel Permeation Chromatography

Table 9.1 summarises the results of the GPC tests on the commercial PA6, the two in

situ masterbatches and the nanocomposites obtained by dilution of the masterbatches

with commercial PA6. The commercial PA6 from Rhodia presents a Mw of about

55000 g/mol, which can be considered as a normal value for this polyamide. The in

situ PA6+20%Sep-NH2 has a molecular weight which is slightly smaller

(Mw=47000) than that of the commercial PA6 but very close to it, while the in situ

PA6+20%Sep masterbatch Mw is only half of it. The reason for such a small value is

221


attributed to a comparatively much higher number of NH2 groups initiating the

polymerisation. It is reminded here that the in situ PA6+20%Sep masterbatch is

initiated by an amount (in mole) of 1-hexadecylamine equivalent to the moles of

primary amines available on the functionalised clays, which instead directly initiated

the polymerisation of the in situ PA6+20%Sep-NH2 masterbatch. It may well be that

1-hexadecylamine is a more efficient initiator than the primary amines attached on

the clay surface and/or that the number of active primary amino groups on the

functionalised sepiolite are overestimated since a part of them might be shielded in

the clay coating layer and be inaccessible to the initiation of ε-caprolactam. The

higher amount of the monofunctional initiator results in an imbalance of end groups

(NH2 and COOH) that have to undergo a condensation reaction, which consequently

limits the molecular weight obtainable in a linear polymer as PA6 [6].

Table 9.1. GPC results of molecular weight averages (Mw and Mn) and

polydispersity (Mw/Mn) of the commercial PA6, the two in situ masterbatches, and

the nanocomposites obtained after dilution of the masterbatches with commercial

PA6. Samples are run in duplicates.

Sample Run No Mw [g/mol] Mn [g/mol] Mw/Mn

Commercial PA6 1 54800 25700 2.1

2 54900 25600 2.1

In situ PA6+20%Sep 3 26700 10800 2.5

4 26200 10400 2.5

In situ PA6+20%Sep-NH2 5 46700 20200 2.3

6 47100 20400 2.3

In situ PA6+3%Sep 7 48100 21900 2.2

8 47700 21800 2.2

In situ PA6+3%Sep-NH2 9 51700 23600 2.2

10 52000 24000 2.2

Nevertheless, the difference in molecular weight between the two masterbatches is

almost annulled after compounding with commercial PA6. For instance the in situ

222


PA6+3%Sep and in situ PA6+3%Sep-NH2 nanocomposites have a very close value

of Mw and the same polydispersity (Mw/Mn). This observation is very important

since it excludes that differences in physical/mechanical properties of in situ

polymerised nanocomposites are caused by different molecular weight and molecular

weight distribution and it allows simple comparisons and discussions.

9.3.2 TGA – Estimation of Inorganic Residue and

Purification of In Situ Masterbatches

As explained earlier, two masterbatches have been produced by in situ

polymerisation. In the case of Sep-NH2, the polymerisation is initiated by the amino

groups attached (or adsorbed) on the clay, from which the polymer chains grow.

100 200 300 400 500 600 700 800 900 1000

20

40

60

80

100

3rd extraction 2nd extraction 1st extraction PA6+20% Sep-NH2

Mas

s [%

]

Temperature [°C]

Figure 9.1. TGA of in situ PA6/Sep-NH2 masterbatch as produced and after three

successive extractions in formic acid (from bottom to top).

223


The amount of covalently bonded PA6 can be roughly estimated by thermo

gravimetric analysis. In Fig. 9.1, the TGA traces of the PA6/SepNH2 masterbatch

subjected to three consecutive washings in formic acid are presented.

Most of the adsorbed, non-covalently linked PA6 chains are probably extracted by

the solvent with the first washing. The amount of inorganic residue stabilises at about

64 wt.%, after the third washing. If it is taken into account that the initial

concentration of inorganic in the PA6/20%Sep-HN2 masterbatch is about 20 %, the

organic grafted to the sepiolite is evaluated in about 11 wt.%. This value comprises

the silane surface functionalisation and the polyamide chains.

The polymerisation of ε-caprolactam to polyamide 6 never goes to completion since

it is an equilibrium reaction [6]. This means that in the course of the process, the

system reaches a point where the concentration of reactants (ε-caprolactam) is in

equilibrium with the concentration of products (polyamide 6) or, in other words, the

rate of the forward reaction is equals to the rate of the reverse reaction.

100 200 300 400 500 600

20

40

60

80

1007 %

Mas

s [%

]

Temperature [°C]

Before Extraction After Extraction

Figure 9.2. TGA in inert atmosphere (N2) of in situ PA6/20%Sep-NH2 masterbatch,

before (solid line) and after (broken line) purification in hot distilled water.

224


An amount of 7-10 wt.% of ε-caprolactam is typically found after the polymerisation

reaction. It is common practise in the nylon industry to extract this unreacted

monomer and reuse it. Taking advantage of the reverse reaction also forms the basis

for the recycling of PA6 by de-polymerisation into ε-caprolactam [7-9]. In this study,

the unreacted monomer in the in situ polymerisation is extracted by hot water and

evaluated by TGA. Fig. 9.2 shows the temperature scans of the PA6/Sep-NH2

masterbatch before and after monomer extraction. In the polymerised compound

before washing, a first step is evident at about 200 °C in the TGA scan, which is

attributed to the degradation/volatilisation of the monomer still present (but also

partially to water). After purification, however, the step almost disappears. This

confirms that the hot water extraction is a simple and viable method to eliminate ε-

caprolactam from the polymer.

In table 9.2, an estimation of the monomer content in the two masterbatches, before

and after purification, is given as a result of the TGA results in oxidative and inert

atmosphere.

Table 9.2. Monomer content in the in situ polymerised masterbatches, PA6/Sep and

PA6/Sep-NH2, before and after extraction in hot water.

Before Purification After Purification

PA6/Sep PA6/Sep-NH2 PA6/Sep PA6/Sep-NH2

N2 Air N2 Air N2 Air N2 Air

Monomer wt. [%]

7.5

8.4

6.7

7.8

2.5

3.2

1.3

1.4

Following the polymerisation and monomer extraction, the two in situ masterbatches

were diluted, by melt compounding with a commercial PA6, to nominal

concentrations of filler of 0.1, 1, 2.5 and 5 wt.%. The real filler content though was

estimated by the inorganic residue after TGA tests. The results are presented in Table

9.3.

225


Table 9.3. Filler content of the different nanocomposites, calculated by the residue of

inorganic phase after TGA tests.

Nominal Filler Content

wt. [%]

TGA Residue (at 900°C)

wt. [%]

In situ PA6+0.1%Sep 0.1 0.75

In situ PA6+1%Sep 1 1.75

In situ PA6+2.5%Sep 2.5 3.35

In situ PA6+5%Sep 5 5.15

In situ PA6+0.1%Sep-NH2 0.1 0.9

In situ PA6+1%Sep-NH2 1 1.7

In situ PA6+2.5%Sep-NH2 2.5 2.95

In situ PA6+5%Sep-NH2 5 5.15


It has already been shown earlier in this thesis how important is the level of

nanofiller dispersion on the final properties of the nanocomposite. The morphology

of the nanocomposites obtained by dilution of the two in situ polymerised

masterbatches is discussed in this paragraph based on SEM and TEM micrographs.

Figure 9.3.a-d, refers to SEM pictures of cold fractured surfaces of the in situ

PA6+5%Sep and in situ PA6+5%Sep-NH2 nanocomposites. In both samples, the

sepiolite (white spots) appears perfectly dispersed in the polymeric matrix (dark

areas). There is a striking difference between the two sets of pictures, though,

specifically the size of the sepiolite particles. It is reminded here that the pictures are

taken from the cross-sectional area of tensile tests specimens. The needle-like

sepiolite clays are sticking out of the plane of the picture in different directions. Most

of the clay particles, in these SEM micrographs, show circular shapes that represent

the projection of the nano-needles in the plane of the cross-section. The bright spots

in the pictures give an approximate but realistic indication of the diameter of the

sepiolite. If Fig.9.3.a is compared with Fig. 6.1.e (Chapter 6), it seems that PA6

226


nanocomposites have a higher filler concentration than PP nanocomposites, although

this is denied by TGA results (Table 9.3). This can be explained by a better

interfacial interaction PA6/Sep rather than PP/Sep. A coating layer of PA6 forms on

sepiolite surface, which makes the sepiolite nanofibres appear bigger and smoother

in PA6/Sep nanocomposites rather than in PP/Sep nanocomposites.

a)

b)

c) d)

Figure 9.3. SEM micrographs of: a)-b) In situ PA6+5%Sep and c)-d) In situ

PA6+5%Sep-NH2, at magnifications respectively of 10000 (left column) and 50000

times (right column).

In the case of in situ PA6/Sep-NH2 (Fig.9.3.d), the diameters of the nano-clays

appear to be more than double the size of those in the in situ PA6/Sep

nanocomposites. This can be explained by the sizing agent of the Sep-NH2 and, more

importantly, by the PA6 chains grafted on the clays.

In Figure 9.4, the TEM micrographs of in situ PA6+5%Sep-NH2 nanocomposites are

also presented. The dispersion of sepiolite is at the level of individual nano-clays, as

already anticipated by the SEM pictures (Fig. 9.3). Moreover, a good interface

227


matrix/filler can be seen with the polymer wetting the sepiolite and forming a low

contact angle.

a) b)

c) d)

Figure 9.4. TEM micrographs of in situ PA6+5%Sep-NH2.

9.3.3 Crystallisation Behaviour

The different nanocomposites based on the two in situ masterbatches were subjected

to DSC tests in order to study the crystallisation and the melting behaviour. The main

results are summarised in Table 9.4. The degree of crystallinity is calculated from the

melting heat (∆Hm), according to the following equation:

cm

c fHH

X ×∆∆

=0

Equation 9.1

Where =240 J/g is the melting enthalpy of 100 % crystalline PA6 [10] and is

a factor that takes in account the excluded mass of inorganic that doesn’t melt. For

0H∆ cf

228


instance, is 0.95 for a 5 wt.% nanocomposite. It can be concluded that there are

no significant variations either in the degree of crystallinity or in the temperature of

crystallisation (T

cf

c) between different concentrations of natural or functionalised

sepiolite.

Table 9.4. Summary of the crystallinity and temperature of crystallisation of the in

situ PA6 nanocomposites.

Sample ∆Hm [J/g] Xc [%] Tc [°C]

PA6 75.6 31 193

PA6+0.1%Sep 75.5 31 193

PA6+1%Sep 75.3 31 193

PA6+2.5%Sep 75.6 31 193

PA6+5%Sep 72.6 30 193

PA6+0.1%Sep-NH2 73.8 31 192

PA6+1%Sep-NH2 75.1 31 193

PA6+2.5%Sep-NH2 77.6 32 193

PA6+5%Sep-NH2 74.5 31 193


The tensile tests of the PA6 nanocomposites, obtained from the two in situ

masterbatches, are displayed in Fig. 9.5 by representative stress-strain curves.

The mechanical properties are summarised in Fig. 9.6-9.9, showing the Young’s

modulus, ultimate tensile stress, strain at break and toughness of the nanocomposites.

As often found in the literature and in this thesis, the presence of sepiolite generally

enhances the stiffness and tensile strength of the polymeric matrix, reducing,

however, the strain at break and the toughness. The stiffness and strength of the in

situ PA6/Sep and in situ PA6/Sep-NH2 nanocomposites are very similar, and the

values slightly diverge only for higher filler content (Fig. 9.6-9.7).

229


a)

0 10 20 30 40 50-10

0102030405060708090

100110120

St

ress

[MPa

]

Strain [%]

PA6 PA6+0.1% Sep PA6+1% Sep PA6+2.5% Sep PA6+5% Sep

b)

0 10 20 30 40 500

10

20

30

40

50

60

70

80

90

100

110

120

Stre

ss [M

Pa]

Strain [%]

PA6 PA6+0.1% Sep-NH2 PA6+1% Sep-NH

2 PA6+2.5% Sep-NH2 PA6+5% Sep-NH2

Figure 9.5. Stress-strain curves of: a) In situ PA6/Sep and b) In situ PA6/Sep-NH2

nanocomposites, at different nominal filler loadings.

230


0 1 2 3 4 52.5

3.0

3.5

4.0

4.5

5.0

5.5

Youn

g's

Mol

udus

[GPa

]

Filler wt. [%]

In-Situ PA6 / Sep In-Situ PA6 / Sep-NH2

Figure 9.6. Elastic Moduli for in situ PA6/Sep and PA6/Sep-NH2 nanocomposites at

different filler loadings.

0 1 2 3 4 575

80

85

90

95

100

105

110

115

UTS

[MPa

]

Filler wt. [%]

In-Situ PA6 / Sep In-Situ PA6 / Sep-NH2

231


Figure 9.7. Ultimate tensile stress for in situ PA6/Sep and PA6/Sep-NH2

nanocomposites at different filler loadings.

Main differences between the two nanocomposites appear, however, in the strain at

break and the toughness. The nanocomposites obtained from the functionalised

sepiolite (Sep-NH2) fails at higher elongations than the nanocomposites from pristine

sepiolite (Sep), as show in Figure 9.8. It is stressed at this point that, although the

data points are quite scattered, the results are statistically valid since they are average

of at least 10 specimens each.

0 1 2 3 4 50

10

20

30

40

50

Stra

in a

t bre

ak [%

]

Filler wt. [%]

In-Situ PA6 / Sep In-Situ PA6 / Sep-NH

2

Figure 9.8. Strain at break for in situ PA6/Sep and PA6/Sep-NH2 nanocomposites at

different filler loadings.

Similar results are provided by Figure 9.9, where the toughness was calculated

simply by integrating the engineering tensile tests curves over the strain, which gives

an indication of the total energy needed to bring the sample to failure. The in situ

PA6/Sep-NH2 nanocomposites appear tougher that the in situ PA6/Sep

nanocomposites in particular at higher concentration of filler, with 100 % increase

relative to 5 wt.% of sepiolite. The reason for this significant difference in ductility

232


and toughness can be related to the nanocomposite structure. As previously

mentioned in the discussion of Fig.9.3, the in situ PA6/Sep-NH2 nanocomposites

present a peculiar morphology, where the sepiolite needles, coated with

propylaminosilane and grafted-from polymerised PA6, are dispersed in the matrix

(commercial PA6). It is suggested that the nano-clay coating can act as a soft

interphase between matrix and filler and hence modify the nanocomposites

deformation mechanism [11-13], without compromising stiffness and strength.

0 1 2 3 4 50

500

1000

1500

Toug

hnes

s (J

m-2)

Filler wt. [%]

In-Situ PA6 / Sep In-Situ PA6 / Sep-NH

2

Figure 9.9. Toughness of in situ PA6 / Sep and PA6 / Sep-NH2 nanocomposites at

different filler loadings, calculated from the integration of the tensile tests curves

force-displacement.

9.3.5 Dynamic Mechanical Analysis

In order to further corroborate the theory of a soft interphase around the

reinforcement phase in the in situ PA6/Sep-NH2, dynamic mechanical analyses

(DMA) are performed on the different nanocomposites. A typical DMA curve of

PA6, over a range of temperature between -140 °C and 150 °C, is provided in Fig.

233


9.10. Generally, the storage modulus decreases with increase of the temperature; the

stiffness dramatically drops from about 3 GPa to about 400 MPa in correspondence

to a relative small interval of temperature (centred in 70 °C).

-150 -100 -50 0 50 100 150

0

1000

2000

3000

4000

5000

6000

Temperature [°C]

Stor

age

Mod

ulus

[MPa

]

0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

γβ

α

Tan δ

Figure 9.10. DMA: storage modulus and tan δ curves of PA6.

This phenomenon is associated with the α-relaxation. Generally, viscoelastic

relaxations correspond to the onset of various types of internal molecular motions.

The α-relaxation is characteristic of the amorphous part of the polymer and

corresponds to the relaxation of cooperatively rearranging regions some tens to

hundreds of repeat units. Such a transition can be easily evidenced by the tanδ curve,

that shows a main peak at about 70 °C, a characteristic temperature also known as

the glass transition temperature (Tg). Phenomenologically, the Tg is the temperature

at which the amorphous part of the polymer experiences the transition from glassy to

rubbery state (and vice versa). The tanδ curve also shows two other minor peaks at

about -50 °C and -120 °C, which are normally referred to the β and γ-relaxation

respectively and which are related to the relaxation of much shorter chain segments

and chain folds [6, 11]. The same relaxations are present in the PA6 nanocomposites.

234


The attention is focused on the α-relaxation. In Figure 9.11, the Tg of the

nanocomposites prepared from the two in situ masterbatches is plotted against

sepiolite loadings.

0 1 2 3 4 564

66

68

70

72

74

76

78

T g [°C

]

Filler Content [%]

In-situ PA6 / Sep In-situ PA6 / Sep-NH2

Figure 9.11. Tg of in situ PA6/Sep and in situ PA6/Sep-NH2 nanocomposites vs.

filler loadings.

The in situ PA6/Sep nanocomposites present glass transition temperatures

statistically invariable with the filler content, if we exclude the effect of extremely

low weight percentages. The Tg of in situ PA6/Sep-NH2 nanocomposites, however,

decreases with increasing amount of sepiolite and reaches values as much as 8 °C

lower than in situ PA6/Sep, for less than 5 wt.% of filler. This behaviour is believed

to be related to the nano-clay coating and to what has already been proposed in the

previous paragraph, as a “soft interphase” between matrix and filler.

It is widely documented in the literature that the Tg of nanocomposites, with well

dispersed nanofillers, can exhibit substantial deviations from the values of the bulk

235


polymer [14-16]. It is generally accepted that this depends on the nature of the

interactions forming between polymer and filler. The Tg should increase in

correspondence to “wetted” polymer/filler interfaces characterised by attractive

interactions (i.e. hydrogen bonds) while the Tg should decrease for weak interfaces.

The two cases above are closely related to two model systems amply studied: thin

polymeric films on a substrate (supported films) with good interfacial interactions,

and thin films at the free surface (free-standing films) [17-20]. The perturbations

induced by the presence of a substrate or a free surface extend for up to 100 nm into

the film, but with different results. At the substrate, interfacial bonds suppress

cooperative segmental mobility, which lead to an increase in Tg, while at the free

surface there is enhanced cooperative segmental mobility and reduced Tg.

In the case of in situ PA6/Sep-NH2 nanocomposites, it is believed that the reduction

of Tg with filler content is induced by the effect of the nanoclay coating that shields

the hydrogen-bond interactions between bulk PA6 and sepiolite, present instead in

the in situ PA6/Sep nanocomposites. This exemplifies the concept of “soft

interphase” polymer/nanofiller proposed previously.

9.4 Conclusions

Two model in situ PA6/sepiolite nanocomposites have been successfully prepared,

differing exclusively in the interfacial interactions: in situ PA6/Sep and in situ

PA6/Sep-NH2. Both nanocomposites have been characterised in terms of

morphology, thermal behaviour, semi-static tensile tests and dynamic mechanical

tests. The morphology, as investigated by SEM, shows a homogeneous dispersion of

the nanoclays into the polar polymeric matrix. Notable enhancements in stiffness and

strength are obtained with only little weight percentage of filler, without modifying

the total amount of crystallinity of the matrix. Interestingly, the grafted-from

nanocomposites (in-situ PA6/Sep-NH2) presented an enhanced strain at break and

toughness (up to 100 % for 5 wt.% of filler) compared with the in situ PA6/Sep

nanocomposites. It is believed that the reason for this difference lie in the modified

236


interface matrix/filler. The in situ PA6/Sep-NH2 comprises what can be described as

a soft interphase, as demonstrated by SEM micrographs and by the reduction in glass

transition temperature, with leads to an enhancement in the energy dissipation.

9.5 References 1. A. Okada, M. Kawasumi, A. Usuki, Y. Kojima, T. Kurauchi, and O.





Research, 1993. 8(5): p. 1179-1184.

3. P. Liu and J. Guo, Polyacrylamide grafted attapulgite (PAM-ATP) via surface-

initiated atom transfer radical polymerization (SI-ATRP) for removal of Hg(II)

ion and dyes. Colloids and Surfaces A: Physicochemical and Engineering

Aspects, 2006. 282: p. 498-503.

4. E. Franchini, Structuration of nano-objects in eposy-based polymer systems :

nanoparitcles and nanoclusters for improved fire retardant properties. 2007,

INSA: Lyon.

5. E. Duquesne, S. Moins, M. Alexandre, and P. Dubois, How can nanohybrids

enhance polyester/sepiolite nanocomposite properties? Macromolecular

Chemistry and Physics, 2007. 208(23): p. 2542-2550.

6. M. Kohan, Nylon Plastics. 1973, New York, London: Wiley-Interscience.

7. JP564550.

8. US3182055.

9. US5869654.



11. J.P. Bell and T. Murayama, Relations between dynamic mechanical properties

and melting behavior of nylon 66 and poly(ethylene terephthalate). Journal of

Polymer Science Part A-2: Polymer Physics, 1969. 7(6): p. 1059-1073.

237


12. W. Jincheng and Y. Gencang, The energy dissipation of particle-reinforced

metal-matrix composite with ductile interphase. Materials Science and

Engineering: A 2001. 303(1): p. 77-81.

13. K. Ding and G.J. Weng, Plasticity of particle-reinforced composites with a

ductile interphase. Journal of Applied Mechanics-Transactions of the Asme,

1998. 65(3): p. 596-604.

14. J.L. Keddie, R.A.L. Jones, and R.A. Cory, Size-Dependent Depression of the

Glass-Transition Temperature in Polymer-Films. Europhysics Letters, 1994.

27(1): p. 59-64.

15. J.A. Forrest, K. Dalnoki-Veress, J.R. Stevens, and J.R. Dutcher, Effect of Free

Surfaces on the Glass Transition Temperature of Thin Polymer Films. Physical

Review Letters, 1996. 77(10): p. 2002-2005.

16. C.J. Ellison and J.M. Torkelson, The Distribution of Glass Transition

Temperatures in Nanoscopically Confined Glass Formers. Nature Materials,

2003. 2: p. 695-700.

17. R.D. Priestley, C.J. Ellison, L.J. Broadbelt, and J.M. Torkelson, Structural

Relaxation of Polymer Glasses at Surfaces, Interfaces, and In Between Science,

2005. 309(5733): p. 456 - 459.

18. P. Rittigstein and J.M. Torkelson, Polymer-nanoparticle interfacial

interactions in polymer nanocomposites: Confinement effects on glass

transition temperature and suppression of physical aging. Journal of Polymer

Science Part B: Polymer Physics, 2006. 44(20): p. 2935 - 2943.

19. M.K. Mundra, S.K. Donthu, V.P. Dravid, and J.M. Torkelson, Effect of Spatial

Confinement on the Glass Transition Temperature of Patterned Polymer

Nanostructures. Nano Letters, 2007. 7: p. 713-718.

20. P. Rittigstein, R.D. Priestley, L.J. Broadbelt, and J.M. Torkelson, Model

Polymer Nanocomposites Provide Understanding of Confinement Effects in

Real Nanocomposites. Nature Materials, 2007. 6: p. 278 - 282.

238

10 Conclusions and Future Work

10.1 Summary

In the last two decades, polymer-clay nanocomposites have attracted great interests

because of the remarkable enhancements in mechanical and physical properties with

minute amount of nano-filler, promising to eliminate the typical compromise

between properties and processability of composite materials. Despite the

expectations created by nano-clays in the academic and industrial communities, their

success has been so far limited. The reasons being (i) the poor dispersion of nano-

clays in polymer matrices, (ii) the often weak interfacial interaction with polymers,

and (iii) the lack of control of nano-clay orientation.

In this thesis, all of the aspects mentioned above have been tackled, studying the

potentials of sepiolite ‘needle-like’ clay as nanofiller in two thermoplastic polymers:

PP and PA6. A particular emphasis is given throughout the thesis on the comparison

of needle-like sepiolite with the more widely investigated ‘platelet-like’ smectite

nanoclays (i.e. montmorillonite MMT).

A full understanding of the characteristics of the constituent phases is essential

before manufacturing any composite materials and to interpret their final properties.

This was exactly the inspiring motto behind Chapter 5, in which sepiolite nanoclay

was thoroughly studied for its morphological, structural, thermal and mechanical

properties. The elementary particles are rigid needles of 650 nm in length and 24 nm

239

CHAPTER 10 – Conclusions and Future Work

in diameter on average, but with a quite broad dimensional distribution. In fact the

mean aspect ratio (L/D) is 27 but it ranges from 10 to 130. Sepiolite is one of the

clays with the highest specific surface area (300 m2g-1, compared with 25-40 m2g-1

for MMT), which is an indicator of its open structure and accessible external surfaces

and its potential ease in dispersing these fibre-like rather than plate-like nanoclays in

polymer matrices. For the first time in literature, the mechanical properties of the

nanoclay have been evaluated via nano-bending tests on individual sepiolite

particles; the elastic modulus of sepiolite was found to be about 200 GPa. With this

preliminary information in place, the reinforcing potential of needle-like (1D)

sepiolite nanofillers was evaluated by means of the shear-lag model and Halpin-Tsai

equations and compared with platelet-like (2D) clays. Theoretically, sepiolite is

expected to give a more effective reinforcement in 1D uniaxially oriented

composites. Vice versa platelet-like clays should provide a higher reinforcement in

the case of 3D random composites.

In Chapter 6, sepiolite is melt compounded in polypropylene. The nanofiller

dispersion in such hydrophobic matrix was unsatisfactory and it was improved by

use of functionalised polymers, as a third phase in the PP/Sep composites, and the

direct functionalisation of the sepiolite surface. Surface modification of the sepiolite

(silanisation) resulted as a very promising route to improve the compatibility

between PP matrix and nanoclays. It guaranteed excellent nanofiller dispersion (up to

1-2 wt.%) and the best mechanical performances among the different systems studied

(in particular for what concerns the yield stress), without the addition of any

compatibiliser. Moreover, in all cases it was interesting to observe that the thermal

degradation of PP, in oxidative atmosphere, is substantially retarded by the presence

of sepiolite and the effect is more evident for increasing loadings of nano-filler, up to

5 wt.%. The mechanical properties of PP/Sep nanocomposites were further analysed

with two micromechanical models (Halpin-Tsai and Pukanszky) and compared with

relevant publications in the scientific literature. Sepiolite seemed to provide a better

reinforcement in terms both of Young’s modulus and yield stress, when compared

with smectite clays with equivalent levels of filler orientation and this was explained

by a better dispersion achievable with needle-like rather than platelet-like clays.

240


The real advantage of needle- or fibre-like nanofillers rather than plate-like

nanoclays is expected to be experienced for 1D composites, as demonstrated in

Chapter 5. Hence highly oriented PP/sepiolite nanocomposite tapes are prepared by

means of solid-state drawing (Chapter7). The orientation of both the polymer

crystallites and the needle-like filler in the tapes was demonstrated by WAXS

studies. Small amounts of sepiolite (< 2.5 wt.%) had an extraordinary effect on the

mechanical properties of PP tapes, for intermediate draw ratios (10 < λ < 20).

However, the DSC analysis suggested that the origin of the improved mechanical

properties was more related to the modification of the semi-crystalline polymer

structure rather than a pure composite reinforcement effect. Whichever the reasons,

sepiolite was demonstrated as an interesting nanofiller for oriented polypropylene

tapes. In fact, beside improved mechanical behaviour, sepiolite clay provides

increased utilisation temperature (lower flowability) and improved thermal and fire

resistance.

PA6/sepiolite nanocomposites were prepared both by melt compounding (Chapter 8)

and in situ polymerisation (Chapter 9). PA6 disperses sepiolite well without any

needs of compatibilisers or clay surface functionalisation. The simple and

environmentally friendly melt-compounding proved as efficient as the in situ

polymerisation for what concerned the dispersion state of nano-filler, making it a

good candidate for industrial and commercial applications. WAXS confirmed the

appearance of γ-phase crystals, against the more stable α-phase, induced by the

presence of sepiolite in PA6. Edge view X-ray tests also established partial

orientation of sepiolite nano-fibres in the longitudinal plane of tensile test specimens.

Notable enhancements in stiffness and strength were obtained with only small

amounts of filler although this was accompanied by a decrease in strain at break. The

mechanical performances were interpreted in terms of two micromechanical models

(Halpin-Tsai and Pukanszky equation) and compared with the results of PA6/MMT

taken from selected literature publications. Sepiolite showed a notable enhancement

in tensile stress compared with MMT, explained by strong hydrogen bonds between

silanols groups on sepiolite surface and the amide groups of the matrix, which are

instead missing in montmorillonite, even after surface functionalisation.

241


Finally, an in situ polymerisation method (Chapter 9) was utilised in particular to

create two PA6/Sep nanocomposite model systems to study the effect of different

interfacial interactions: in situ PA6/Sep and in situ PA6/Sep-NH2. The former is a

simple polymerisation of PA6 “in presence of” natural sepiolite, while the last is a

polymerisation “grafted from” the functionalised surface of sepiolite (Sep-NH2).

Interestingly, the “grafted from” nanocomposites (in situ PA6/Sep-NH2) presented an

enhanced strain at break and toughness (up to 100 % for 5 wt.% of filler) compared

with the in situ PA6/Sep nanocomposites, without compromising Young’s modulus

and tensile strength. It is believed that the reason for this difference laid in the

modified interface matrix/filler. The in situ PA6/Sep-NH2 showed what can be

described as a soft interphase, as demonstrated by SEM micrographs and by the

reduction in glass transition temperature, with leads to an enhancement in energy

dissipation.

10.2 Future Work

From both a theoretical and experimental point of view, this thesis indicates that

sepiolite is a promising nanofiller, particularly for unidirectionally oriented (1D)

nanocomposite, in comparison with the more investigated smectite clays (i.e. MMT).

The main reasons are the peculiar needle-like shape (instead of the plate-like shape),

the relative ease in dispersing sepiolite in polymeric matrices and the presence of

silanol groups on the sepiolite surface which allows viable and effective

functionalisation, necessary for hydrophobic polymers as PP, or directly strong

interactions with hydrophilic polymers as PA6. It seems natural to focus future works

principally in the area of unidirectional polymer/sepiolite nanocomposites. PP/Sep

nanocomposites tapes have already shown interesting physical properties (Chapter 7)

that can find commercial applications. These can be in the area of synthetic textile

fibres (mainly constituted by polyolefin, polyamide and polyester fibres) or in the

niche but more profitable market of all-PP composites such as Curv® or PURE®.

242


All-PP are self-reinforced composite materials, in which bulk isotropic PP (matrix) is

reinforced by highly oriented and anisotropic polypropylene films, tapes or fibres [1-

3]. Such composites are characterised by high stiffness, low weight, high impact

strength and 100 % recyclability, and have already applications in automotive,

construction and anti-ballistic fields, among others [4]. Sepiolite could be used to

reinforce the highly oriented PP phase while at the same time improve the thermal

and fire resistance, and barrier properties. This strategy can be seen as providing

another hierarchical level to the material engineering at the nanometre scale, without

adding more complexity to the manufacturing of the composite. In fact it could be

easily implemented just by using PP masterbatches loaded with small amounts of

sepiolite clay, on established manufacturing line.

Very interesting results can be also expected from oriented nanocomposites based on

polyamide or polyester matrices. In fact, strong interfacial interaction between matrix

and filler and excellent dispersion of sepiolite in PA6 were already observed in

Chapter 8, without the need of any compatibilisers or surface modification.

Another way to produce unidirectional nanocomposites is via the technique of

electrospinning. The electrospun polymers and composites reported in literature are

already numerous, including nanoclay composites [5-8]. To the best of the author’s

knowledge sepiolite has never been attempted to be included in electrospun fibres.

Figure 10.1 presents some initial work on fibres of a blend of HPC and PEO filled

with sepiolite, successfully prepared via electrospinning. The TEM micrograph in

Fig.10.1.a-c show small bundles of sepiolite in a HPC/PEO fibre (100-150 nm in

diameter). It is also possible to organise single electrospun fibres into a given fibrous

structure, by simply changing the fibre collecting set-up. Fig.10.1.d and Fig.10.1.e,

for instance, show how to achieve aligned fibres or how these can be twisted into

yarns.

243


a) b)

c) d)

e)

Figure 10.1. TEM micrographs of: a)-c) HPC/PEO electrospun fibres filled with

sepiolite needle-like clay, d) aligned electrospun fibres and e) electrospun fibres

twisted into a yarn. The electrospun fibres in the Figures 10.1.a-c are about 200nm

in diameter, which gives a scale for those micrographs.

Beside the search for applications of sepiolite clay in the area of oriented

nanocomposites, fundamental research is still to be carried out and basic questions to

be answered.

244


One is certainly the effect of melt compounding on the aspect ratio of the sepiolite

itself. In Chapter 6, breakage of the nano-needles was already observed. For the

future a more systematic study of the effect of processing on the aspect ratio of the

sepiolite is to be carried out.

Also, the beneficial effect of sepiolite on the thermal resistance of PP was shown via

TGA tests (Chapter 6), suggesting a fire-retardant effect in such a matrix. However

the most recognised apparatus for fire tests is the Cone Calorimeter. Such tests will

be performed in the future on sepiolite nanocomposites.

Finally more nano-mechanical tests can be performed. For the first time the Young’s

modulus of nanoclay has been measured. Future work will be focus on the ultimate

tensile strength and the strain at break of sepiolite. Another interesting research

objective will be to evaluate the strength of the interphase polymer/clay by

specifically designed pull-out tests. In analogy with some work on carbon nanotubes

[9, 10], sepiolite could be attached to a SPM tip, embedded in a liquid resin and

pulled out from it after solidification or curing.

10.3 References

1. J. Loos, T. Schimanski, J. Hofman, T. Peijs, and P.J. Lemstra, Morphological

investigations of polypropylene single-fibre reinforced polypropylene model

composites Polymer, 2001. 42(8): p. 3827-3834.

2. B. Alcock, N.O. Cabrera, N.-M. Barkoula, J. Loos, and T. Peijs, The

mechanical properties of unidirectional all-polypropylene composites.

Composites, Part A: Applied Science and Manufacturing, 2006. 37(5): p. 716-

726.

3. B. Alcock, N.O. Cabrera, N.-M. Barkoula, C.T. Reynolds, L.E. Govaert, and T.

Peijs, The effect of temperature and strain rate on the mechanical properties of

highly oriented polypropylene tapes and all-polypropylene composites.

Composites Science and Technology, 2007. 67(10): p. 2061-2070.

4. http://www.pure-composites.com.

245


5. W.E. Teo and S. Ramakrishna, A review on electrospinning design and

nanofibre assemblies. Nanotechnology, 2006. 17: p. R89-R106.

6. Z.-M. Huang, Y.-Z. Zhangb, M. Kotakic, and S. Ramakrishna, A review on

polymer nanofibers by electrospinning and their applications in

nanocomposites. Composites Science and Technology, 2003. 63(15): p. 2223-

2253.

7. L. Li, L.M. Bellan, H.G. Craighead, and M.W. Frey, Formation and properties

of nylon-6 and nylon-6/montmorillonite composite nanofibers. Polymer, 2006.

47: p. 6208-6217.

8. V.K. Daga, M.E. Helgeson, and N.J. Wagner, Electrospinning of Neat and

Laponite-Filled Aqueous Poly(ethylene oxide) Solutions. Journal of Polymer

Science, Part B: Polymer Physics, 2006. 44: p. 1608–1617.

9. A.H. Barber, S.R. Cohen, and H.D. Wagner, Measurement of carbon nanotube-

polymer interfacial strength. Applied Physics Letters, 2003. 82(23): p. 4140-

4142.

10. A.H. Barber, S.R. Cohen, A. Eitan, L.S. Schadler, and H.D. Wagner, Fracture

transitions at a carbon-nanotube/polymer interface. Advanced Materials, 2006.

18(1): p. 83-87.

246

List of Author’s Publications [1-7] 1. E. Bilotti, H. Deng, R. Zhang, D. Lu, W. Bras, H.R. Fischer, and T. Peijs,

Highly oriented multifunctional polypropylene / sepiolite clay nanocomposite

tapes. Submitted to Polymer.

2. E. Bilotti, H. Deng, R. Zhang, D. Lu, W. Bras, H.R. Fischer, and T. Peijs, PA6

nanocomposites based on needle-like sepiolite nanoclay. To be published.

3. H. Deng, T. Skipa, R. Zhang, D. Lellinger, E. Bilotti, I. Alig, and T. Peis, Effect

of melting and crystallisation on the conductive network in conductive polymer

composite. Submitted to Polymer.

4. H. Deng, R. Zhang, E. Bilotti, and T. Peijs, Effective reinforcement of carbon

nanotubes in polypropylene matrices. Submitted to Journal of Applied Polymer

Science.

5. H. Deng, R. Zhang, E. Bilotti, J. Loos, and T. Peijs, Conductive polymer tape

containing highly oriented carbon nanofillers. Accepted in Journal of Applied

Polymer Science, 15 October. DOI 10.1002 / app.29624.

6. E. Bilotti, H.R. Fischer, and T. Peijs, Polymer nanocomposites based on

needle-like sepiolite clays: Effect of functionalized polymers on the dispersion

of nanofiller, crystallinity, and mechanical properties. Journal of Applied

Polymer Science. 107(2): p. 1116-1123.

7. J. Ma, E. Bilotti, T. Peijs, and J.A. Darr, Preparation of polypropylene/sepiolite

nanocomposites using supercritical CO2 assisted mixing. European Polymer

Journal. 43(12): p. 4931-4939.

247

thesis bilotti 2009

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