Modification of carbon fiber / epoxy matrix interphase in...

N° d’ordre 2014ISAL00127 Année 2014

Thèse

Modification of carbon fiber / epoxy matrix interphase in a

composite material: design of a self-healing interphase

by introducing thermally reversible Diels-Alder adducts

présentée devant

L’Institut National des Sciences Appliquées de Lyon

pour obtenir

Le Diplôme de docteur

École Doctorale: Matériaux de lyon

Spécialité: Matériaux Composites

par

Wenyong ZHANG Soutenue le 11-12-2014 devant la Commission d’Examen:

JURY Reviewer Daniel Hanoch WAGNER Professor

(Weizmann Institute of Science, Israel) Reviewer Michel NARDIN Directeur de recherche CNRS (Institut de Science des Matériaux de Mulhouse, France) Examiner Philippe ZINCK Professor (Université des Science et Technologies de Lille, France) Supervisor Jean François GÉRARD Professor

(INSA de LYON, France) Co-Supervisor Jannick DUCHET-RUMEAU Professor (INSA de LYON, France)

INSA de Lyon-Université de Lyon

Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2014ISAL0127/these.pdf © [W. Zhang], [2014], INSA de Lyon, tous droits réservés

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Remerciements Ces travaux ont été financé par le Conseil de Bourse Chinoise (CSC) et effectués

au sein du l’UMR CNRS 5223/Ingénierie des Matériaux Polymères, au Laboratoire

de Matériaux Macromoléculaires à l’Institut National des Sciences Appliquées de

Lyon.

Tout d’abord, je tiens à exprimer ma plus sincère gratitude au Professeur

Jean François GERARD, directeur de la recherche INSA de Lyon, et Professeur

Jannick DUCHET-RUMEAU pour m’avoir accueilli dans leur laboratoire et dirigé

pendant ces quatre années. Sans leurs précieux conseils, leurs aides efficaces et

aussi leurs encouragements, je n’aurais pas fini ces travaux.

Je tiens ensuite à remercier Daniel Hanoch WAGNER, Professeur à L’Institut

Weizmann des Sciences, et Michel NARDIN, Professeur à L’Institut de Science

des Matériaux de Mulhouse pour avoir accepté d’être les rapporteurs de cette

thèse. De plus, je remercie également Philippe ZINCK, Professeur à L’Université

des Science et Technologies de Lille, examinateur, pour sa présence et sa

contribution à mon jury de thèse.

J’aimerais aussi remercier Laurence MASSIN, ingénieur à l’Institut de

Recherches sur la Catalyse et l'Environnement de LYON, pour avoir réalisé les

analyses XPS. Je remercie Fernande BOISSON, ingénieur de recherche du

laboratoire IMP, et Cécile CHAMIGNON, ingénieur d’étude du laboratoire IMP,

pour la réalisation de l’étude RMN. Je voudrais particulièrement remercie Guilhem

QUINTARD, Technicien du laboratoire IMP, pour son aide précieuse durant les

tests mécaniques. Merci aussi à Georgiy SMOLYAKOV, post-doctorant au

laboratoire IMP, qui m’a formé à l’AFM. Je remercie toutes les personnes qui j’ai

rencontré durant ma thèse et qui m’ont aidé à la réalisation de ce travail.


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Mes remerciements vont également à mes aimables co-bureaux, Marie, Céline,

Elie, Thomas, Noelie, Krisd, Nicolas et Nour qui m’ont accompagné durant ces

quatre belles années passées à l’IMP. Je remercie également mes amis chinois,

Senbin, Zhen, Jing (il), Xibo, Biao et Jing (elle) pour leur amitié et leurs aides sans

réservations. Merci à tous les camarades avec qui ce fût en plaisir de travaille.

Enfin, un énorme remerciement à ma famille, mes parents, mon frère et

particulièrement à ma épouse qui m’a accompagné pendant les périodes difficiles.


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Abstract

A thermally self-healable carbon/epoxy interphase was designed based on

Diels-Alder (D-A) thermally reversible covalent bonds. The D-A modified

interphase was formed between maleimide groups grafted on carbon fiber surface

and furan groups introduced into epoxy network. The self-healing ability was

characterized by a micromechanical approach using the micro-droplet debonding

test. In this work, carbon fiber surface underwent a three-step treatment to graft

maleimide groups, including HNO3 oxidization, tetraethylenepentamine (TEPA)

amination, and bismaleimide (BMI) grafting. The fiber surface physico-chemical

modifications after each treatment step were characterized by microscopies (SEM,

and AFM) and spectroscopies (XPS, and ATR-FTIR). The matrix modification was

carried on mixing furfuryl glycidyl ether (FGE) into epoxy/amine network and the

properties of modified matrix were studied by TGA, DSC, ATR-FTIR, and tensile

tests. The reversible character of Diels-Alder bond was also followed by DSC, TGA,

and NMR. The interfacial mechanical properties and the self-healing abilities of the

D-A modified interphases, built by combining DGEBA-FGE/amine matrix with a

serial of BMI-grafted carbon fibers tuned as a function of the oxidization time were

investigated. At last, since FGE plays a double-role in D-A modified interfacial

system, i.e. chain extender in epoxy network and self-healing agent in the

interphase, the influences of FGE content in matrix on the mechanical properties of

interphase and also on the mechanical properties of cured matrix were evaluated.

As a consequence, this study allowed to achieve the best process to build a

thermally self-healable carbon/epoxy interphase based on thermally reversible

Diels-Alder covalent bonds. The formed interphase has not only the successive


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self-healable abilities but also the required mechanical properties. Additionally, the

overall mechanical properties of the composite material based on this interphase

will not be weakened significantly after the interfacial modifications.

Thesis title: Modification of carbon fiber / epoxy matrix interphase in a composite

material: Design of a self-healing interphase by introducing thermally reversible

Diels-Alder adducts

KEY WORDS: Diels-Alder; Interface/Interphase; Self-healing; Carbon Fiber;

Epoxy; Composites; Micro-droplet; Debonding.

Résumé

Une interphase fibre de carbone/matrice époxy thermiquement auto-réparable a

été construite sur la base de liaisons covalentes Diels-Alder (D-A) thermiquement

réversibles. L’interphase modifiée par D-A a été formée en greffant des groupes

maléimide sur la surface de la fibre de carbone et en introduisant des groupes

furane dans le réseau polyépoxy. La capacité d’auto-réparation interfaciale a été

caractérisée par le test de déchaussement de la micro-goutte. La surface de la

fibre de carbone a subi un traitement en trois étapes : (i) oxydation par l’acide

nitrique, (ii) amination par la tétraéthylènepentamine (TEPA) et (iii) greffage de

bismaléimide (BMI). Après chaque étape de traitement, les modifications

physico-chimiques de surfaces de la fibre ont été caractérisées par microscopies

(MEB et AFM) et par spectroscopies (XPS, et ATR-FTIR). La modification de la

matrice a été effectuée en copolymérisant le furfuryl glycidyl éther (FGE) au


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réseau époxy/amine et les propriétés de la matrice ont été évaluées par TGA,

DSC, ATR-FTIR, et traction uniaxiale. Le caractère réversible des liaisons

Diels-Alder a été également vérifié par DSC, TGA et RMN. Pour caractériser les

capacités d'auto-réparation de l’interphase modifiée par D-A, les propriétés

mécaniques et les capacités d'auto-réparation de l'interphase construite en

combinant la matrice DGEBA-FGE/amine avec une série de fibres de carbone

greffés par BMI ont été mesurées en fonction du temps d’oxydation préalable au

greffage (gouvernant la réactivité de la fibre de carbone). Enfin, car le FGE joue un

double rôle dans le système interfacial modifié par D-A, à la fois dans l’architecture

en intervenant comme allongeur de chaîne entre nœuds de réticulation du réseau

époxyde et au niveau de l’interphase en contribuant dans la formation des liaisons

réversibles, l'influence de la concentration de FGE dans la matrice a été étudiée

sur les propriétés mécaniques de l'interphase et également sur les propriétés

mécaniques de la matrice. Par conséquent, ce travail a permis d’aboutir à la

procédure optimale pour construire une interphase fibre de carbone/époxy

thermiquement auto-réparable basée sur des liaisons covalentes Diels-Alder

(thermo réversibles). L'interphase ainsi formée possède non seulement des

capacités d’auto-réparations multiples, mais également des propriétés

mécaniques compatibles avec une approche ‘matériau composite’. En effet, les

propriétés mécaniques globales des matériaux composites, comme attendu, sont

dépendantes des caractéristiques de cette interphase mais ne seront pas réduites

par la présence de celle-ci notamment pour assurer la durabilité du matériau

composite.


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Titre de thèse: Modification de l’interphase du matériau composite fibre de

carbone /matrice époxyde: Design d’une interphase auto-réparable basée sur des

liaisons Diels-Alder thermiquement réversibles

Mot clés: Liaisons Diels-Alder; Interface/ Interphase; Auto-réparabilité; Fibre de

Carbone; Epoxy; Composite; Déchaussement de la Micro-goutte


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Table of content:

General introduction

Part I Bibliography A. Interphase of carbon fiber reinforced epoxy matrix based composite . 7

A.1 Structure of carbon/epoxy interphase ............................................................... 7 A.1.1 Dimensions of interphase .................................................................................. 7

A.1.2 Interfacial adhesion: Basis ................................................................................ 8

A.1.3 Specific features of the interphase in case of ex-PAN carbon fiber ................. 10 A.2 Modification of carbon/epoxy interphase ........................................................ 11

A.2.1 Aims of interfacial modifications in carbon fiber/epoxy systems....................... 11

A.2.2 Modification of carbon fiber surface ................................................................. 15

A.2.2.1 Oxidative treatment .......................................................................................... 15

A.2.2.2 Non-oxidative treatments ................................................................................. 18

A.2.2.3 Fiber surface characterization .......................................................................... 24

A.2.3 Modification of epoxy matrix ............................................................................ 25

B. Characterization of interfacial mechanical properties of composite materials ......................................................................................................... 27

B.1 Introduction ..................................................................................................... 27 B.2 Micron-scale methods .................................................................................... 28

B.2.1 Single fiber pull-out test ................................................................................... 28 B.2.1.1 Introduction ....................................................................................................... 28

B.2.1.2 Parameters influencing the single fiber pull-out test ........................................ 32

B.2.1.3 Interfacial failure modes ................................................................................... 37

B.2.1.4 Interfacial failure criteria ................................................................................... 38

B.2.2 Single fiber fragmentation test ......................................................................... 42

B.2.3 Single fiber push-out test or micro-indentation test .......................................... 43

B.2.4 Compression test ............................................................................................ 45 B.3 Conclusion ...................................................................................................... 46

C. Self-healable composite materials .......................................................... 47 C.1 Extrinsic self-healing for polymer materials .................................................... 48

C.1.1 Embedded microcapsules ............................................................................... 49

C.1.2 Embedded hollow fibers .................................................................................. 51

C.1.3 Tri-dimensional microvascular network ........................................................... 52 C.2 Intrinsic self-healing for polymer materials ..................................................... 55

C.2.1 Self-healing based on thermoplastic polymer .................................................. 56


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C.2.2 Self-healing process based on reversible bonds ............................................. 57 C.2.2.1 Reversible covalent bonds ............................................................................... 57

C.2.2.2 Supramolecular assemblies ............................................................................. 61

C.2.2.3 Metal ligands .................................................................................................... 64

C.3 Selection of a strategy for interfacial self-healing ........................................... 66 D. Conclusion ................................................................................................ 71

Part II Design and characterization of a self-healable interphase A. Methodology to generate a self-healable carbon/epoxy interphase ... 75

B. Carbon fiber surface treatments and characterizations ....................... 76 B.1 Introduction ..................................................................................................... 76 B.2 Carbon fiber surface treatment ....................................................................... 78 B.3 Characterization of carbon fiber at different treatment steps .......................... 80

B.3.1 Characterization of carbon fiber after oxidization step ..................................... 81

B.3.1.1 Selection of initial carbon fiber ......................................................................... 81

B.3.1.2 Fiber surface topography after oxidization step ............................................... 83

B.3.1.3 Carbon fiber surface chemical composition after oxidization step ................... 85

B.3.1.4 Influence of HNO3 oxidization on fiber mechanical properties ......................... 87

B.3.2 Characterization of carbon fiber after grafting with TEPA and BMI .................. 89 B.3.2.1 Selection of amination agent ............................................................................ 89

B.3.2.2 Fiber surface topography after grafting with TEPA and BMI ........................... 90

B.3.2.3 Fiber surface chemical composition after grafting with TEPA and BMI ........... 91

B.3.3 Conclusion ...................................................................................................... 94

C. Epoxy matrix composition and characterization ................................... 95 C.1 Composition and curing of matrix system ....................................................... 95 C.2 Matrix characterization ................................................................................... 97

C.2.1 Chemical structure of fully cured furan-functionalized epoxy networks ............ 97

C.2.2 Physical properties of furan functionalized epoxy network .............................. 98

D. Formation and evaluation of thermally reversible self-healable interphase .................................................................................................... 100

D.1 Investigation of Diels-Alder reaction ............................................................. 100 D.2 Formation of interphase containing thermo-reversible bonds ....................... 105 D.3 Interphase micromechanics: micro-droplet debonding test .......................... 108

D.3.1 Specimen preparation ................................................................................... 108


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D.3.2 Process of test .............................................................................................. 109

E. Conclusion .............................................................................................. 112

Part III

Interfacial mechanical properties and self-healing ability of conventional carbon/epoxy interphase A. Introduction ............................................................................................. 117

B. Sample geometrical parameters for micro-droplet debonding test .. 118 B.1 Inter-relationship and significance of embedded length (Le), droplet diameter (Dd), and contact angles (θ and ϕ) ...................................................................... 118 B.2 Significance of free length (Lf) ...................................................................... 122 B.3 Statistics of fiber diameter (Df) ...................................................................... 124

C. Interfacial mechanical properties ......................................................... 124 C.1 Interfacial fracture mode analysis ................................................................. 124 C.2 Analysis of debonding process(es)............................................................... 126

C.2.1 Analysis of the first debonding force ............................................................. 126

C.2.2 Definition of the embedded length (Le) range ................................................ 129

C.2.3 Analysis of the second debonding force ........................................................ 130

C.2.4 Analysis of the frictional force ....................................................................... 131 C.3 Determination of the interfacial shear strength (IFSS).................................. 132

C.3.1 Interfacial shear strength measured during the first debonding (IFSS1) ......... 132

C.3.2 Interfacial shear strength measured during the second debonding (IFSS2) ... 134

C.3.3 Evaluation of the interfacial shear strength of the four reference interphases 135

D. Evaluation of interfacial self-healing ability ......................................... 136

E. Conclusion .............................................................................................. 139

Part IV

Interfacial mechanical properties and self-healing ability of D-A modified interphase A. Influence of carbon fiber surface treatment on interfacial mechanical performances and interfacial self-healing ability of the D-A modified interphase .................................................................................................... 145

A.1 Introduction ................................................................................................... 145 A.2 Wettability of furan-modified epoxy matrix on different fiber surfaces .......... 147


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A.3 Analysis of interfacial shear strength ............................................................ 148 A.4 Evaluation of interfacial self-healing ability ................................................... 150

A.4.1 Feasibility of interfacial self-healing ............................................................... 150

A.4.2 Interfacial self-healing efficiency.................................................................... 155

A.4.3 Comparison healing abilities of the different D-A modified interphases ......... 157

A.4.4 Repeatability of interfacial self-healing .......................................................... 160 A.5 Conclusion .................................................................................................... 162

B. Influence of FGE content on D-A modified interphase properties .... 164 B.1 Introduction ................................................................................................... 164 B.2 Influence of FGE content on the interfacial mechanical properties ............... 165

B.2.1 Influence on the interfacial bonding strength ................................................. 165

B.2.2 Influence of FGE content on the interfacial toughness .................................. 166 B.3 Influence of the FGE content on the interfacial self-healing efficiency.......... 168

B.3.1 Characterization of the interfacial self-healing efficiency based on IFSS ....... 168

B.3.2 Characterization of interfacial self-healing efficiency based on Gic ................ 170 B.4 Conclusion .................................................................................................... 173

General conclusions and perspectives ................. 175

Annexe ...................................................................... 181

Reference ................................................................. 187


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General introduction

A composite material is defined as a material that is constituted with two phases:

matrix phase and reinforcing phase. The chemical or/and physical properties of

these two phases are usually different and must be assort with each other. The

combination of them should produce a synergy effect and lead to some new

properties or even bring a new function to the material. Any material coincident

with above conditions at the same time can be considered as composite material.

The excellent properties of composite materials such as high specific

strength, high specific modulus, and relatively high toughness combined with

simple processes make them widely used in some industrial fields. For example, in

the aeronautics and aerospace fields, the structural material should be not only

strong enough to ensure the security but also light enough to reduce the transport

cost. However, for the conventional materials (aluminium-based materials), the

safety issues are proportional to the weight. As a consequence, composite

materials, especially the carbon fiber reinforced advanced composite materials,

get more and more interests in these fields. Furthermore, the applications of

composite materials in the fields of civil industries have also developed

significantly such as in automobile industry and sports equipment. Nowadays, the

environmental pollution caused by automobile has been more and more serious.

Using polymer based composite materials to replace conventional metallic

material can effectively reduce the dead weight of cars and thus decreasing the

consumption of fossil fuels. As the development of composite materials, its

application range is broadening gradually and almost involves all the industries.

A composite material is formed by joining two different phases together by an


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area, named interface. However, the researchers have confirmed [Drzal1983,

Wu2014] that an interface of a composite material is not two-dimensional (2D), but

a narrow region formed by interpenetration between matrix and reinforcing

material (and/or its surface treatment) at a sub-micronic or nanometric scale. The

structural, physical, chemical, and mechanical properties of this region are totally

different to the neat constituent ones. For this reason, this area is further named as

‘interphase’ but not interface. In a composite material, the interphase always plays

a very important role. For example, as a carbon fiber-based composite material

undergoes a sudden impact, it will instantaneously appear a huge number of

micro-cracks mainly in the interphase region. The appearance of these

micro-cracks can absorb a large part of impact energy and can improve the

toughness of the material. In fact, the interphase is a transition layer where stress

concentration and defects happen. In the early stage of loading of a composite

material, the existing defects and stress concentration will trigger a large number

of imperceptible micro-cracks in the region of interphase. If we do nothing for these

cracks, they will continue to propagate in the process of use and extend inside

matrix until the whole material is broken. This broken process is unavoidable for all

composite materials. For any type of loading conditions (tensile, compression,

flexural, shearing, and torsion), the load applied on the matrix will be transmitted

through interphase to the reinforcing phase, thereby the mechanical properties of

the material will be improved.

The aim of this PhD thesis is to design a carbon fiber and epoxy based

composite material whose interphase has a thermal self-healing ability. In other

words, the micro-cracks in interphase of this type of composite material would be

healed by heating, i.e. the life of use can be extended and catastrophic failure of


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the composite material can be avoided. The realization of this subject is based on

building a carbon/epoxy interphase from conventional thermally reversible

Diels-Alder (D-A) bonds resulting from the reaction between furan and maleimide

groups. As micro-cracks appear in the interphase region, a subsequent heating will

allow the D-A adducts on the corresponding surfaces of the cracks to re-establish

between crack lips. Moreover, this healing process can be repeated many times.

Part I summarizes some conventional methods used to evidence interfaces

and main features about carbon fiber/epoxy interphase modification, interphase

mechanical characterization, and self-healing applied to polymer and composite

materials. After this part, the most relevant methodology for designing a

self-healable interphase as well as the characterization methods will also be

proposed.

Part II details the protocols for modifying carbon fiber surface and matrix that

will considered for designing self-healable interphase. The micro-droplet

debonding test will be introduced for characterizing interfacial mechanical

properties and self-healing efficiency.

Part III is dedicated to the characterization of the interfacial mechanical

properties of conventional carbon/epoxy interphases (non-D-A modified, i.e.

considered as ‘references’) using the micro-droplet debonding test and the

relevant sample geometrical parameters will be optimized. The mechanical

properties and self-healing abilities of these reference interphases will be also

discussed.

Part IV reports the discussions of experimental results of micro-mechanical

test on interphases designed to be self-healable. The feasibility of Diels-Alder

reaction based self-healing interphase is confirmed and the best conditions of fiber


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surface treatment and matrix modification are determined. In this part, the

interfacial mechanical properties of this self-healing interphase are also discussed.

Finally, the conclusions are summarized in Part V and some perspectives

are discussed.


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Part I Bibliography A. Interphase of carbon fiber reinforced epoxy matrix based composite . 7

A.1 Structure of carbon/epoxy interphase ............................................................... 7 A.1.1 Dimensions of interphase .................................................................................. 7

A.1.2 Interfacial adhesion: Basis ................................................................................ 8

A.1.3 Specific features of the interphase in case of ex-PAN carbon fiber ................. 10 A.2 Modification of carbon/epoxy interphase ........................................................ 11

A.2.1 Aims of interfacial modifications in carbon fiber/epoxy systems....................... 11

A.2.2 Modification of carbon fiber surface ................................................................. 15

A.2.2.1 Oxidative treatment .......................................................................................... 15

A.2.2.2 Non-oxidative treatments ................................................................................. 18

A.2.2.3 Fiber surface characterization .......................................................................... 24

A.2.3 Modification of epoxy matrix ............................................................................ 25

B. Characterization of interfacial mechanical properties of composite materials ......................................................................................................... 27

B.1 Introduction ..................................................................................................... 27 B.2 Micron-scale methods .................................................................................... 28

B.2.1 Single fiber pull-out test ................................................................................... 28 B.2.1.1 Introduction ....................................................................................................... 28

B.2.1.2 Parameters influencing the single fiber pull-out test ........................................ 32

B.2.1.3 Interfacial failure modes ................................................................................... 37

B.2.1.4 Interfacial failure criteria ................................................................................... 38

B.2.2 Single fiber fragmentation test ......................................................................... 42

B.2.3 Single fiber push-out test or micro-indentation test .......................................... 43

B.2.4 Compression test ............................................................................................ 45 B.3 Conclusion ...................................................................................................... 46

C. Self-healable composite materials .......................................................... 47 C.1 Extrinsic self-healing for polymer materials .................................................... 48

C.1.1 Embedded microcapsules ............................................................................... 49

C.1.2 Embedded hollow fibers .................................................................................. 51

C.1.3 Tri-dimensional microvascular network ........................................................... 52 C.2 Intrinsic self-healing for polymer materials ..................................................... 55

C.2.1 Self-healing based on thermoplastic polymer .................................................. 56

C.2.2 Self-healing process based on reversible bonds ............................................. 57

C.2.2.1 Reversible covalent bonds ............................................................................... 57

C.2.2.2 Supramolecular assemblies ............................................................................. 61

C.2.2.3 Metal ligands .................................................................................................... 64 C.3 Selection of a strategy for interfacial self-healing ........................................... 66


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D. Conclusion ................................................................................................ 71


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Part I Bibliography A. Interphase of carbon fiber reinforced epoxy matrix based composite

A.1 Structure of carbon/epoxy interphase

The interphase formed between a carbon fiber and an epoxy matrix displays

specific features. Its nature is not a simple contact surface between carbon fiber

and epoxy matrix, but a interpenetration and ‘connection’ of each of the two

phases from physical or/and chemical effects. The interpenetration and

‘connection’ of each phase not only makes interphase have a certain thickness, but

also leads to generate different physic-chemical properties to the interphase. This

section will describe in detail the structure of interphase formed between carbon

fiber and epoxy resin, including the dimensions of interphase, the interfacial

adhesion, and some limitations of interphase.

A.1.1 Dimensions of interphase

The earliest definition of interphase can be dated back to the year of 1974

[Metcalfe1974]. The researchers, at that time, considered interface as a layer with

certain thickness but not a mathematical plane (2D). However, they did not know

the exact structure of interphase because of the limit of characterization

techniques. In the early 1990s, Guigon et al. [Guigon1991, Guigon1992,

Guigon1994] firstly characterized the micro-structure of interphase in carbon fiber

reinforced composite by using high resolution transmission electron microscopy

(HRTEM). According to the TEM images, the existence of the interface was

confirmed to be a layer with a certain thickness which is resulted from


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inter-diffusion and interaction between the fiber surface and the polymer matrix.

However, their research works did not indicated the precise dimensions of

interphase. In 2013, Wu et al. [Wu2014] measured the thickness of T300/epoxy

interphase by TEM. For getting an accurate result, the ion etching technique was

used to treat sample surface. The sliced sample was observed under various

directions of incident electron beam. According to the phase-contrast, the position

and micro-structure of the interphase were identified and the interphase thickness

was evaluated to be about 200 nm (Fig. I-1).

Figure I-1. TEM image of T300/epoxy interphase [Wu2014].

A.1.2 Interfacial adhesion: Basis

The bonding mode of interphase is an extremely complex theme, which is

governed by the physicochemical properties of fiber surface and matrix. Generally,

the formation of interphase is considered as the result of two effects: chemical

effect and physical effect [Kinloch1980, Wake1978, Kim1998]. The ‘chemical

effect’ means that the active functional groups on fiber surface react with the


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compatible chemical groups of matrix to form chemical bonds, by which the fiber

surface and matrix are combined together. The ability of the interphase to transfer

stress from the polymer matrix to the fiber depends on the number and bond

energy of the bonds existing at the interphase. About ‘physical effect’ based

adhesion theory, four acceptable origins could be considered [Kinloch1980]:

mechanical interlocking, diffusion, electrostatics, and adsorption. The mechanical

interlocking theory states that the polymer chains could enter and anchor in the

graphite layers at the fiber surface (cavities and micro-grooves) or tangles with

surface polymer chains (for coated fiber), by which fiber surface and matrix are

interlocked. The effect of interlocking is determined by fiber surface topography

and the polymer chain mobility during contact, i.e. processing. The diffusion

theory proposes that the molecules of two different phases diffuse into each other.

Obviously, this theory can’t be considered for explaining the formation of an

interphase between carbon fiber and a polymer, except if the fiber surface was

coated previously. In the later case, the inter-diffusion occurs between the polymer

chains and the components of the coating/sizing layer. The electrostatic theory

states that the electrostatic charges between fiber and matrix are different and the

interaction of positive and negative charges (electrostatic force) can lead to

interfacial bonding which depends on the charge density. Although electrostatic

force is usually tiny between fiber and matrix, it can be strengthened when the fiber

surface is treated with some specific coupling agents. The absorption theory is

the most generally accepted one. The involved adhesion forces are so-called

secondary forces, including Van der Waals’ force and hydrogen bonds. In general,

the secondary forces could be considered as non-efficient for interfacial bonding.

In fact, for the most of interphases encountered in fiber based composite materials,


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the interfacial adhesion strength can’t be explained with only one of the theories

mentioned above. One needs to consider all of them depending on the chemistry

of the fiber surface, topography, and the physicochemical properties of the matrix.

A.1.3 Specific features of the interphase in case of ex-PAN carbon fiber

It’s well-know that carbon fiber has a specific turbostratic structure which is

composed by crystallized graphite planes, as schematically shown in Figure I-2.

These disordered carbon layers are covered by a thin skin of few well-organized

basal planes [Guigon1994]. This special structure makes the carbon fiber surface

be very inert and only few functional groups are present on it. Thus, it’s impossible

to form a large number of chemical bonds between carbon fiber without any

surface treatment and epoxy matrix. Moreover, the surface energy of graphite

plane is extremely low, which makes absorption theory irrelevant on a carbon fiber

surface. For these reasons, the interface of composite material reinforced by

carbon fibers without any treatment will be very weak. That is the reason for which

there are so many investigators interested in carbon fiber surface modification to

promote interfacial adhesion.

Figure I-2. Schematic drawing of a three-dimensional model of a carbon fiber [Bennett1978].


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In a composite material, the function of interphase, besides combining fiber

surface and polymer matrix, is to transfer stress from the matrix phase to

reinforcing fiber leading to reinforcement. However, the mechanical properties of

reinforcing material and matrix are extremely different, especially for carbon fiber

and epoxy matrix. These obvious differences will lead to stress concentration close

to the interface, i.e. in the interphase. In other words, when a composite material is

loaded, the interphase could be damaged before other regions of the composite

material.

The first specific feature described above related to an intrinsic limitation of

carbon fiber to establish required interfacial interactions which can be improved by

fiber surface modification. However, the second feature which exists in all kinds of

composite materials, i.e. enhanced stress application to the interphase, can’t be

avoided but can be optimized from interphase design.

A.2 Modification of carbon/epoxy interphase

Interphase is a transition region formed by the fiber surface and the matrix. For

optimizing its structure and properties, the structure and properties of fiber surface

and matrix should be considered independently or simultaneously.

A.2.1 Aims of interfacial modifications in carbon fiber/epoxy systems

Carbon fiber is a relevant reinforcing material for composite due to its high specific

strength and modulus. Nevertheless, the inert character of its graphite-like surface

limits the adhesion with polymer matrices such as epoxy network. Thus, the most

important objective for interfacial modification is the improvement of the

interphase region as it plays an essential role for transferring stress from the


12

polymer matrix to the fiber and it’s the place for stress concentration. As a

consequence, the elastic behavior and toughness of the interphase are essential.

As a conclusion, the interface/phase modification is motivated according to the

following requirements:

a) Interface/phase strengthening

Improved interfacial properties can be achieved from the generation of interfacial

chemical bonds and/or physical effects. The majority of fiber surface treatments

offer to introduce active functional groups on fiber surface, especially after

oxidization and grafting treatments. Some of these groups can react with the

reactive functional groups of the matrix and form interfacial chemical bonds. These

bonds can dramatically increase the interfacial strength. For example, Peng et al.

[Peng2013] introduced a large amount of primary and secondary amine groups on

carbon fiber surface by means of grafting method. As the treated fibers are

combined with epoxy matrix to form carbon fiber reinforced composite materials,

these amine groups on fiber surface act as a curing agent. The authors highlighted

that the interfacial shear strength (IFSS) after treatment displays an increase of 85%

compared to that without treatment (Fig. I-3). Furthermore, almost all the active

functional groups introduced by surface treatment are polar groups. For a polar

epoxy matrix, these polar groups on the fiber surface can undoubtedly promote the

wettability of matrix, phenomenon very important for interfacial adhesion.


13

Figure I-3. Evolution of the interfacial shear strength (IFSS) as a function of the concentration of

poly(amidoamine) (PAMAM) [Peng2013].

Some treatments methods can also etch fiber surface, as the weak layer and

attached particles on fiber surface are removed leading to a rougher fiber surface.

The rough surface is favorable for physical anchoring of matrix increasing the

interfacial adhesion strength. When the etching effect is strong enough, some

micro grooves or cavities are formed on fiber surface (Fig. I-4). Some of scientists

[Rashkovan1997, Meng2013, Yu2014] considered these grooves and cavities as

surface defects which will affect mechanical properties of carbon fiber and also

mechanical performances of the resulting composite material. However, others

[Zhang2004, Li2008] indicated that these ‘defects’ allowed matrix chain segments

to enter leading to a good mechanical interlocking with matrix, which is an efficient

way to get high stress transfer levels. Connecting (by grafting or CVD) carbon

nanotubes on fiber surface to form a multi-scale reinforced composite material is a

relatively new subject. But its interfacial strengthening principle is not very clear

nowadays.


14

Figure I-4. AFM images of treated and untreated ex-PAN carbon fiber surface [Meng2013].

b) Interphase toughening

Hitherto, few works deal with the interphase toughening. A simple concept for

interphase toughening is to introduce some flexible polymer chains between fiber

and matrix. Varelidis et al. [Varelidis1999] have studied the interlaminar fracture

toughness of carbon fiber reinforced composite material, whose reinforcing carbon

fibers were coated with polyamide 66 (Nylon66). The results indicated that the

flexible chains of Nylon66 existing in interphase surely increased the interfacial

fracture toughness. In this system, there is not chemical bond between Nylon66

and fiber surface. So it’s impossible to get an ideal interfacial strength. Zhao et al.

[Zhao2010] have grafted a layer of uniform octaglycidyldimethylsilyl POSS on

carbon fiber surface. This flexible layer of silicone increased impact toughness

from 2.62 to 3.59 J. Simultaneously, the interlaminar shear strength (ILSS) has

also been increased from 68.8 to 90.5 MPa, which can be attributed to the covalent

bonds formed between POSS and fiber surface (Fig. I-5).


15

Figure I-5. Mechanical properties of the composite materials based on untreated and POSS-grafted

carbon fibers. a) ILSS and b) impact toughness [Zhao2010].

A.2.2 Modification of carbon fiber surface

Carbon fiber surface treatment can be generally classified into two types: oxidative

and non-oxidative treatments. According to oxidative medium, oxidative treatments

are further divided into dry oxidization and wet oxidization [Kim1998].

A.2.2.1 Oxidative treatment

a) Wet oxidization

Wet oxidization treatments are achieved in liquid oxidizing agents, such as nitric

acid [PittemanJr1997, Wu1995, Serp1998], hydrogen peroxide [Meng2013,

Guo2009], supercritical water [Bai2010], and potassium persulfate [Yu2014] etc.

For increasing the oxidization efficiency, these oxidizing agents generally can be

used together with some additives. For example, Bai et al. [Bai2010] investigated


16

the oxidization effect of supercritical water on carbon fiber surface. They found that

it had a slight ability to oxidize carbon fiber surface. However, when they added

oxygen into the supercritical water, the oxidization ability increased strongly. In

addition, the researchers have also found that the carbon fiber surface topography

might be changed by etching effect of the oxidization agent [Yu2014]. For example,

the pits and/or grooves appear on fiber surface. These pits and/or grooves can

even damage the original mechanical properties of carbon fiber when the number

or the deepness of them reaches a certain value. The change of topography

depends on oxidization conditions, including the type of oxidizing agent,

oxidization temperature, and exposure time. For example, both Pittmanjr

[PittmanJr1997] and Serp [Serp1998] have tried to treat the carbon fiber surface by

nitric acid with a same concentration. The treatment conditions of Pittmanjr are

115°C for different times from 20min to 90min, while the Serp’s conditions are 80°C

for 300min. The different temperatures and durations of oxidization treatments

lead to almost opposite conclusions. Pittmanjr indicated that the amount and

dimension of pits on fiber surface increased with oxidization time, whereas Serp

[Serp1998] concluded that the fiber surface morphology was not affected

significantly for the considered conditions.

The other wet oxidization treatment which proceeds faster and uniformly is

commonly used for commercial carbon fibers. It’s denoted as electrolytic or anodic

oxidization [Park2000]. This method proceeds with considering the carbon fiber as

anode in an electrolytic reaction with different electrolyte, for example sulphuric

acid aqueous, ammonium bicarbonate, or sodium hydroxide [Kozlowski1984]. This

electrolytic treatment introduces active functional groups on surface and

dramatically changes the topography of fiber surface.


17

b) Dry oxidization

Dry oxidization treatment is carried out in atmosphere of gas. The gas can be

some oxygen containing gases such as oxygen [Boudou2003], ozone [Li2008],

carbon dioxide [Serp1998], and air [Fukunaga1999, Zhang2004, Lee1997,

Xie1990] etc. Although the most part of these gases are oxidative at room

temperature, their oxidizing power is not sufficient for treating the extremely inert

carbon fiber surface and this treatment process always needs to be performed at a

relatively high temperature. Zhang et al. [Zhang2004] have oxidized carbon fiber

surface under air atmosphere with varying temperature, included between 450 and

600 °C during 1 hour. They have shown that, at 450 °C, fiber surface morphology

does not change in a large way compared with the untreated fiber. While, when the

temperature increases to 500 °C, the fiber surface becomes relatively rough and

some pieces of tiny fragments appear on fiber surface. At 600 °C, the oxidization

becomes more aggressive and a large number of micro-pits could be generated on

the fiber surface. Obviously, the carbon fiber surface treatment in air with high

temperature has a negative effect on structure of fibers, and therefore damaging

their mechanical properties. In fact, no matter the nature of the gas, the strong

oxidizing character under high temperature of these gases will etch fiber surface

significantly. The plasma oxidization is another relatively mild treatment which will

not etch fiber surface significantly [Boudou2003]. This treatment is apt to turn

oxygen containing gases into plasma. The electric particles (ion and free electron)

in plasma atmosphere will attack fiber surface and will provide a large amount of

oxygen functional groups on fiber surface. Although this method seems more

suitable for carbon fiber surface treatment, it needs a special plasma reactor.

No matter which oxidization methods are adopted, all of them have a same


18

consequence: the amount of oxygen-containing functional groups on fiber surface

will be more or less increased. When the treatment conditions are strong enough,

the fiber surface might be etched and even lead to lower mechanical properties of

fibers. Figure I-6 displays the possible functional groups existing on fiber surface

after oxidization treatments.

OO C

OC

C

O

OH OH O

O

O

OHO

O

O

O

Carboxyl Phenolic Hydroxyl Quinone-type Carbonyl

Normal Lactone Fluorescene type lactone

Carboxylic Acid Anhydride

Cyclic Peroxide

Figure I-6. Generated functional groups on fiber surface after oxidization treatment of ex-PAN

carbon fibers [Jones1991].

A.2.2.2 Non-oxidative treatments

The non-oxidative treatments do not change the chemical structure of fiber surface

itself, but could introduce some additional reactive species close to the fiber

surface. This method will not influence the mechanical properties of carbon fiber

and generally can effectively improve the interfacial properties.


19

a) Coating/sizing

The coating consists in forming an extremely thin polymer layer on the fiber

surface, which is usually considered for commercial carbon fibers. The coating

could be either applied from a solution or an emulsion of polymer(s) and additive(s)

[Guigon1994, Cheng1994, Paipetis1996, Broyles1998, Dilsiz1999, Dilsiz2000,

Dai2011, Zhang2012]. The coating mainly has two functions [Paipetis1996,

Dilsiz2000]: 1) protection of fiber surface during fiber handling and re-processing

as well as composite manufacturing; 2) promotion of fiber surface wettability and

further increase of the interfacial adhesion. The first function is accepted almost by

all the researchers in composite field. However, the second objective could be

considered according to different ways. Dilsiz et al. [Dilsiz1999, Dilsiz2000] studied

the physic-chemistry of the fibers coated with polyetherimide and poly (thiol-

arylene phosphine oxide). They mentioned that coating can induce a decrease of

both polar and dispersive components and surface functional groups concentration.

While, Dai et al. [Dai2011] proposed an opposite conclusion by studying

commercially coated carbon fibers (T300B and T700SC), i.e. coated fiber surface

have a higher concentration of activated carbon atoms and a higher polar surface

energy than that of non-coated fiber. It means that the type of coating agent

directly determines the physicochemical nature of the fiber surface which is

contact with the polymer matrix. Additionally, the second function of coating layer

heavily depends on the compatibility between coating layer and matrix

[Cheng1994, Guigon1994]. For example, Yumitori et al. [Yumitori1994] have

reported that the carbon fiber coated with brominated epoxy resin can effectively

improve the interphase with the polyethersulfone (PES) matrix, whereas it leads to

a weak interphase in the case of epoxy matrix system. For PES matrix, strong


20

dipolar-dipolar interactions take place between bromine atoms from the fiber

coating and PES, leading to interpenetration between coating layer and matrix. But

for epoxy matrix, this kind of interaction does not exist and only physical absorption

occurs between sizing layer and matrix (Fig. I-7). Besides these parameters, the

state of coating components before deposition (polymer solution or emulsion)

[Cheng1994], the coating layer thickness [Gerard1988], and the molecular weight

of coating components [Zhang2012] also play a key role.

Figure I-7. Interphase structures between Br-epoxy coated carbon fiber and two different polymer

matrices [Yumitori1994].

b) Surface grafting

The surface grafting is one of the most effective methods to activate and/or

functionalize carbon fiber surface. The grafted materials can be multifunctional

polymer chains, and covalently bonded carbon nanotubes, etc. For example,

Pittmanjr et al. [PittmanJr1997a] grafted long-chain tetraethylenepentamine (TEPA)

onto nitric acid oxidized ex-PAN carbon fiber surface, resulting in a dramatic

increase of the amount of amine groups on fiber surface (Fig. I-8). These amine

groups can significantly increase the wettability of fiber surface and can be reacted

during curing in the case of epoxy matrix based composite material. For


21

introducing more active amine groups on fiber surface, poly(amido amine)

(PAMAM) dendrimer that contains more -NH2 and -NH- has been also used

[Peng2013, Mei2010].

Figure I-8. Proposed mechanisms of grafting tetraethelynepentamine (TEPA) onto oxidized carbon

fiber surface [PittermanJr1997a].

Carbon nanotubes (CNT) could be also grafted on the carbon fiber surface to

design high performance multi-scale reinforced composite materials [He2007,

Laachachi2008]. For example, Laachachi et al. [Laachachi2008] have treated CNT

and CF by acid oxidization and thermal treatment, respectively. These protocols

allowed to form covalent bonds between CNT and CF, i.e. to attach the CNTs onto

carbon fiber. Figure I-9 simply demonstrates this grafting process. For increasing

the grafting efficiency of CNT, Mei et al. [Mei2010] have also grafted PAMAM on

oxidized carbon fiber surface in order to increase the quantity of active points

(-NH2 and -NH-) that can react with oxidized CNT.

Figure I-9. Chemical bonding of functionalized nanotubes on carbon fibers from subsequent

esterification, anhydridation or amidization of carboxyl groups formed during oxidization treatment

[Laachachi2008].


22

All the grafting methods mentioned above are based on oxidized carbon fibers.

Actually, some grafting reactions on carbon fiber surface can be carried out directly

without any prior treatment. In Section A.1, we mentioned that initial carbon fiber

(without any treatment) was covered by a thin skin of a few well-organized

graphite-like planes. However, these graphite crystalline planes are not perfect and

numbers of defects exist. Close to the defects and at the edge of planes, C=C

double bonds could be identified, as well as some double bonds even conjugated

[Stein1985]. The double bonds on carbon fiber surface provide a possibility to graft

polymer chains directly on its surface. Barbier et al. [Barbier1990] have

electrochemically oxidized ethylene-diamine, by which radical cations were

produced. These radical cations can attack fiber surface double bonds to form C-N

covalent bonds, thereby grafting ethylene-diamine on fiber surface. Moreover, this

grafting method can be used for a wide variety of molecules which offers a great

scope of possible surface modification. Peng [Peng1995] and Buttry [Buttry1999]

have indicated that the double bonds on carbon fiber surface can be considered as

electrophile allowing the reaction with nucleophilic amine groups following a

Michael-like addition. This reaction allows to graft any long-chain molecules

containing amine groups on fiber surface both at room and elevated temperatures.

Subsequently, Severini et al. [Severini2002] have also proposed to graft maleic

anhydride or tetracyanoethylene on conjugated C=C bonds on the fiber surface by

Diels-Alder addition. Although direct grafting reaction on inert carbon fiber surface

is possible, the grafting efficiency is significantly limited by the number of defects

on fiber surface. As a consequence, these surface treatment methods were not

reported many times in the literature.


23

c) Others

The above parts introduced the conventional non-oxidative methods for carbon

fiber surface. In fact, other specific non-oxidative methods have been mentioned

somewhere, such as chemical vapor deposition (CVD). Zhu et al. [Zhu2003]

synthesized carbon nanotube on carbon fiber surface by using CVD. The gas for

nanotube synthesis was a mixture of methane and hydrogen with different

contents of methane. Iron nanoparticles were deposited on fiber surface as

catalyst. By considering this method, carbon nanotubes successfully grew on the

fiber surface with a relatively low reaction temperature (650-800°C). Zhao et al.

[Zhao2005] reported that CVD with floating catalyst method is more favorable for

the growth of CNT on carbon fibers compared to the conventional thermal CVD

and the morphology of new formed nanotube significantly depended on the

reaction temperature. Low temperature deposition leads to produce short

MWCNTs with larger diameter while high temperature allows to make long

MWCNTs with a smaller diameter. The hyper-thermal and hypo-thermal treatments

have also been checked by some investigators. Wang [Wang2006] and Dai

[Dai2011a] heated carbon fibers under vacuum at different high temperatures.

They found that high temperature could improve the thermal stability and the

degree of graphitization of carbon fiber and decrease the amount of functional

groups on the surface. This treatment could be achieved not only under vacuum

but also under some inert atmosphere [Lee1997]. The hypo-thermal treatment

[Rashkovan1997, Zhang2004] of carbon fibers could occur in some liquid gas

media which are at extremely low temperature, for example, liquid carbon dioxide,

nitrogen, helium, or their vapors. After treatment, the weak layer and attached

particles on fiber surface will be removed and fiber surface becomes rougher.


24

A.2.2.3 Fiber surface characterization

After treatment, both the surface topography and surface chemical structure will be

changed on a small scale. How to characterize these changes has become one of

the most popular subjects in the field of carbon materials. The development of new

modern microscopy techniques allow to observe fiber surface at nanometer scale

or even at the scale of a few atoms, such as transmission electronic microscopy

(TEM) and atomic force microscopy (AFM). Other techniques, for example gas

absorption, and Raman microprobe spectrometry (RMS) etc. can also be used to

analyze fiber surface quantitatively. About the surface chemical composition, X-ray

photoelectron spectroscopy (XPS), Attenuated Total Reflectance Fourier

Transform Infrared spectroscopy (ATR-FTIR), and inverse gas chromatography

(IGC) are conventional methods that can be used for qualitatively or

semi-quantitatively analyzing the fiber surface. Although acid and base uptake can

estimate the content of acidic and basic functional groups on fiber surface, its

precision is affected by many factors. Table I-1 lays out some well-known methods

for carbon fiber surface characterization.

Technique Characteristic Reference

Scanning electron microscope

(SEM)

Direct observation of the fiber

surface with high magnification

Morita1986, Smiley1993,

Boudou2003

Atomic force microscope

(AFM)

3D surface topography of fiber

and measurement of surface

roughness

Smiley1993

Gas absorption Determination of fiber surface

area and pore volume Boudou2003

Raman microprobe spectrometry

(RMS)

Characterization of the surface

regularity of fiber with about 100

nm analysis depth

Boudou2003


25

Inverse gas chromatography (IGC)

Determination of dispersive part

of surface free energy (γsd) and

acidity and basicity constants of

fiber surface

Vickers2000, Lindsay2007,

Narjes2007

X-ray photoelectron spectroscopy

(XPS)

Quantification or

semi-quantification of surface

functional groups concentration

Kozlowski1984, Wang2000,

Bradley1993, Smiley1993,

Gardner1995, Gardner1996,

Vickers2000, Pamula2003

Lindsay2007

Ion-scattering spectroscopy

(ISS)

Determination of mass spectrum

of the outermost one or two

atomic layers on the carbon fiber

surfaces. Be used as a

complementarity of XPS

Gardner1995, Gardner1996

Attenuated total reflectance Fourier

transform infrared spectroscopy

(ATR-FTIR)

Quantitative analysis of the

functional groups on fiber

surface

Sellitti1990, Ohwaki1993,

Pamula2003

Acid and base uptake

Determination of the content of

acidic and basic functional

groups on fiber surface

Pamula2003, PittmanJr2004

Solid state Nuclear Magnetic

Resonance (CP-MAS NMR)

Analysis of the chemical

composition of fiber surface with

a relatively low resolution

Morita1986

Electron energy loss spectroscopy

(EELS)

Determination of the atomic

content of low atomic mass

elements (C, N, O) on fiber

surface

Morita1986, Pamula2003

Table I-1. The methods for carbon fiber surface characterization.

A.2.3 Modification of epoxy matrix

For improving interfacial performance of carbon fiber reinforced composite

material, the researchers have paid large attention to treat the inert carbon fiber

surface. However, there are few works which considered the direct design of the

interfacial region, i.e. the interphase, from the modification of matrix close to the


26

fiber surface in order to improve interfacial performances. In fact, the interfacial

improvement can be also achieved by matrix modification. He et al. [He2010]

compared the interlaminar shear strength (ILSS) of carbon/epoxy composites as a

function of epoxy nature. They reported that the more epoxy groups existed in the

epoxy prepolymer, the better the ILSS was. Hence it can be predicted that the

bonding strength between fiber and matrix can be achieved by introducing more

functional groups on prepolymer of matrix by chemical modification. These

functional groups should be able to react with the functional groups on fiber

surface or physically interact with fiber surface. Unfortunately, very few works

deals with that. Another matrix modification tends to add a second reinforcing

phase to improve the interfacial properties. For example, Jiang et al. [Jiang2008]

added a certain ratio of fullerene nanoparticles in epoxy matrix to improve the

interfacial properties of carbon fiber reinforced composite material. The interfacial

bond strength between carbon fiber and epoxy matrix was seriously enhanced.

This method can be considered not only as matrix modification to ameliorate the

interfacial properties but also as the constitution of a multi-scale reinforced

composite material. Additionally, the interfacial performance can be improved by

changing the physical state of low viscosity matrix. For example, Huang et al.

[Huang2002] treated liquid matrix by ultrasound before making a carbon fiber

reinforced composite material. The interfacial properties were improved and they

attributed this phenomenon to ultrasonic treatment that has mixed matrix more

uniformly and decreased its viscosity.


27

B. Characterization of interfacial mechanical properties of composite materials

B.1 Introduction

The interfacial modification methods mentioned-above are mainly applied to

improve the mechanical properties of interphase, for example interfacial strength

and toughness, thereby obtaining an optimized composite material. For studying

the effect of these interfacial modifications, relevant characterization methods

should be considered. Until now, the methods for interfacial characterization can

be classified into two types according to the scale of measurement, including

micron-scale and macro-scale. The micron-scale methods are generally carried

out with a micro-composite which is composed by a single fiber (or a certain

number of fibers in some cases) and a low quantity of matrix, such as single fiber

pull-out, single fiber indentation, and single fiber fragmentation, etc. The

characteristics of the mechanical behavior of interphase obtained experimentally

using such methods, include interfacial shear strength (IFSS) and critical energy

release rate (Gic). These methods generally possess relatively definite failure

modes and accurate analysis models, which will be detailed later. The

macro-scale methods, including short beam shear, Iosipescu’s shear test, [±45°]

tensile test, and [10°] off-axis tensile test, etc., use composites as unidirectional

and 0/90° laminates composite materials to enhance as much as possible the

interphase contribution and lead to the fracture mechanisms related to interface

such as interfacial debonding. The samples used here are more similar as a real

composite so that the testing results are more instructive. However, the fracture

mode inside the samples is usually very complex, which involves not only

interfacial debonding but also fiber breakage and matrix fracture. Obviously, for our


28

purpose (interphase self-healing) which only concerns about the performances of

interphase, the micron-scale methods were selected.

B.2 Micron-scale methods

B.2.1 Single fiber pull-out test

B.2.1.1 Introduction

The single fiber pull-out test involves embedding one end of a single fiber within

matrix (can be performed both with a thermoplastic or a thermoset matrix) with the

fiber axis perpendicular to the surface. The force which is required to pull out fiber

from matrix can be used to evaluate the mechanical properties of interphase. The

evaluation methods will be introduced in the section B.2.1.4. Figure I-10 displays

the schematic of single fiber pull-out test.

Figure I-10. Schematic illustration of single fiber full-out test. The matrix droplet is fixed on substrate

[Zhandarov2005].

Although this method can be used to directly measure interfacial bond strength of

composite, there are some inherent limitations [Miller1987]. Generally, the fibers

used for composing a composite material are very thin (~10 μm). If we want a

successful pulling out, the fracture tensile strength of the fiber between matrix

droplet and clamp (used for clamping fiber’s free end) should be smaller than the

interfacial strength which proposes a critical embedded length Lecri, shown as


29

follow:

𝑳𝒆𝒆𝒆𝒆 = 𝝈𝒇𝑫𝒇𝟒𝝉𝒖𝒖𝒖

(1),

where σf is the fracture tensile strength of fiber, Df is the fiber diameter and τult is

the interfacial shear strength. When the embedded length is longer than this critical

length, the fiber will be broken before pulling out. Since the tensile strength of

single fiber (σf) decreases as a function of fiber length [Pardini2002], the

relationship mentioned above can be used only in the case with a constant fiber

length between matrix droplet and clamp to estimate the critical value. In fact, the

value of critical embedded length can be changed with the variation of fiber length

between matrix droplet and clamp. For the commonly used composite systems

(glass/epoxy, carbon/epoxy, or Kevlar/epoxy, etc.), the value of Lecri is very small,

which makes the positioning of the fiber, i.e. the control of the embedded length

extremely difficult. In addition, for matrices which offer a good wetting of the fiber

surface, a meniscus could be formed and render the determination of the

embedded length even more difficult. For solving these problems, Miller et al.

[Miller1987, Gaur1989] have proposed an improved method, the micro-droplet

debonding test which involves the deposition of a small quantity of liquid matrix

onto the surface of a single fiber. After solidification, there is an ellipsoidal droplet

formed on fiber (Fig. I-11a). The testing needs a special clamp by which the droplet

is blocked between two fixed parallel blades (microvise) and the upper fiber end is

connected with a sensor of dynamometer (Fig. I-11b). When the upper fiber end is

pulled upward, the droplet will be debonded.


30

(a) (b) Figure I-11. Conventional micro-droplet (a) and the protocol for micro-droplet debonding test (b)

[Miller1987].

Some researchers [Day1998, Ash2003, Choi2009, Kang2009] have studied the

stress state mapping of micro-droplet debonding test by finite element analysis

(FEA). They indicated that, during debonding test, the stress distribution along the

embedded length is not uniform, but decreases gradually from the point where

fiber enters matrix. In other words, at the upper end of the droplet, an important

stress concentration occurs which will significantly decrease the determination

accuracy of interfacial shear strength (IFSS) as the method proceeds from the use

of micro-droplets having different dimensions. The finite element analysis results

also showed that the stress concentration can be weakened by changing the

shape of droplet. Figure I-12 displays some new sample shapes of debonding test

which can effectively decrease the stress concentration at upper end of the

micro-droplet.


31

(a) (b)

Figure I-12. New sample shapes for micro-debonding test: (a) quasi-disc type sample [Choi2009];

(b) cylinder sample [Morlin2012].

It should be noticed that all the specimens used in above methods are composed

with a single fiber. The interactions between neighboring fibers are completely

ignored [Qiu1993]. Additionally, Rao et al. [Rao1991] indicated that the mechanical

properties of epoxy matrix could vary with droplet size as the concentration of

curing agent could be different. In order to consider the interphase in an

environment close to the one in composite materials, some methods using several

fibers were developed such as multi-fiber pull out test and micro-bundle pull out

test (Fig. I-13). However, these methods have not been widely applied because of

the great difficulty in sample preparation.

(a) (b)

Figure I-13. Schematic illustrations of multi-fiber pull out test (a) and micro-bundle pull out test (b)

[Rao1991].


32

B.2.1.2 Parameters influencing the single fiber pull-out test

In theory, IFSS is an intrinsic parameter of interphase, which should be a constant

value for a given fiber/matrix system. However, the experimental results obtained

for a given fiber/matrix system from the different methods reported previously are

generally not in agreement. There are many factors influencing the results such as

the shape of matrix micro-droplet, and the geometrical parameters of microvise,

etc. [Gaur1989, Wagner1993, Day1998, Subramani2012]. This section

summarizes the influence of some parameters which were analyzed for using

single fiber pull-out test in an appropriate manner.

The shape of matrix micro-droplet

The shape of matrix micro-droplet, including overall shape and dimension of

micro-droplet, is one of the most important factors for single fiber pull-out test. Choi

et al. [Choi2009, Kang2009] analyzed the stress state for different micro-droplet

shapes (droplet, sphere, quasi-disc and cylinder) by finite element analysis (FEA).

The results indicated that the stress concentration occurs mainly in the meniscus

region. In other words, the shape with a small meniscus should be chosen to

decrease the stress concentration. However, controlling the micro-droplet shape is

not easy to be achieved, especially on a micron-scale. For this reason, the

ellipsoidal droplet formed without intervention has been widely applied. In fact,

micro-droplet shape is only one of the numerous factors which influence the size of

meniscus. Its size is more determined by the wettability of matrix on fiber surface.

For some fiber/matrix interphases which result from a poor or moderate wettability

such as carbon/epoxy interphase, the meniscus size remains very small and can

be ignored [Gaur1989].


33

Figure I-14. Dependence of shear strength with embedded length for the first micro-droplets

(shearing force imposed by microvise) and the second micro-droplets (shearing force imposed by

already sheared first droplet) [Gaur1989].

The micro-droplet dimensions, such as embedded length and micro-droplet

diameter (only for micro-droplet debonding test), are also very important factors for

single fiber pull-out test. The works of Gaur and Miller [Gaur1989] indicated that

IFSS decreases somewhat with the increasing of embedded area (or embedded

length) (Fig. I-14). The authors related this dependence to stress concentration.

When a fiber is pulled out from matrix, the largest part of the load is applied to

initiate and propagate cracks. As the embedded area increase, the part of the

applied stress concentrated at the top end of the droplet increases, which will lead

to earlier fracture, i.e. a lower apparent IFSS. In fact, this hypothesis is only

qualitative. More appropriate analysis can be considered from the point of view of

fracture criteria which will be introduced later. Furthermore, for a defined interfacial

system, the successful debonding requires that the embedded length should be

included in a limited range. The upper limit has been already presented and

denoted as the critical embedded length (Lecri). The lower limit for droplet has been


34

mentioned at first by Scheer and Nairn [Scheer1995]. They have demonstrated

that, for a certain fiber/matrix system, there is a strict linear relationship between

embedded length and droplet diameter. Moreover, the diameter of droplet can’t be

smaller than the diameter of fiber that is constant. Hence there exists a minimum

embedded length, which is determined by the fiber diameter and the relationship

between embedded length and droplet diameter.

The configuration of micro-vise

Haaksma and Cehelnik [Haaksma1989] indicated that the configuration of

microvise has a significant effect on IFSS measured by micro-droplet debonding

test. However, Wagner et al. [Wagner1993] who compared three different

configurations of microvise obtained an opposite conclusion (Fig. I-15), i.e. IFSS is

insensitive to the type of loading configuration. They concluded that the different

results are provided from different stiffness of matrix.

Figure I-15. Different microvise configurations of micro-droplet debonding system: (a) parallel

knife-edge loading (two-point loading), (b) parallel cone plate loading (two-point loading), and (c)

circular loading [Wagner1993].

Chou et al. [Chou1994] studied the effect of microvise gap width on micro-droplet

debonding test. The results indicated that increasing microvise gap width leads to

an overestimation of intrinsic IFSS. They attributed this phenomenon to the

distribution of applied force in axial and radial directions (Fig. I-16(a)). Owing to the


35

special shape of micro-droplet (ellipsoid), the force applied to the contact point (A

or B) can be divided into two directions, axial and radial. By increasing the

microvise gap width, the repartition between these two directions will be changed,

i.e. the axial force will decrease and the radial force will increase. The axial force is

considered as the main source of shear force in interphase which leads to

interfacial debonding. The decreasing of the proportion of axial force in applied

force will surely lead to an overestimation of true value of IFSS. In addition, the

increasing of radial force (compressive force) will increase the gripping force of the

fiber, resulting in a higher debonding force. For these reasons, Chou proposed that

the gap width should be chosen as small as practicable to get the best evaluation.

(a)

(b)

Figure I-16. (a) Force analysis for the micro-droplet debonding test. A and B are contact points

[Chou1994]; (b) Demonstration of vise angle [Heilhecker2000].

The dependence of IFSS as a function of microvise gap width has also been


36

demonstrated by Ash et al. [Ash2003] with finite element analysis (FEA) method.

The only difference is that Chou has chosen gap width as researching object,

whereas Ash has chosen vise angle, represented in Figure I-16(b). In fact, these

two parameters have the same geometric signification for an ellipsoidal droplet,

thus can be considered as one parameter for micro-droplet debonding test.

Others

The parameters mentioned-above have been widely studied. Besides these, there

are other parameters which influence more or less the result of debonding test

such as the loading rate and the skill of operator etc. Gaur and Miler [Gaur1989]

studied the relationship between loading rate and IFSS of micro-droplet debonding

test and showed that the measured IFSS increased with a great increasing of

loading rate (more than two orders of magnitude). The operator’s skill has also

been considered as an influencing factor for the micro-droplet debonding test

[Wagner1993, Morlin2012]. In addition, the samples for pull-out test are very

sensible to the environmental factors, especially to temperature and humidity, thus,

during the testing, these factors should also be controlled.

(a) (b)

Figure I-17. Conventional force-displacement traces during single fiber pull-out test with a stable

crack propagation (a) [Zhandarov2000] and with an unstable crack propagation (b) [Kang2009]. Fd

and Fs are debonding force and frictional force, respectively.


37

B.2.1.3 Interfacial failure modes

For the different fiber/matrix systems, the interfacial failure modes are different,

mainly determined by the mechanical properties of interphase. Zhandarov and

coworkers [Zhandarov2000, 2002, 2005, 2006; Pisanova2001, 2001a] have done

a large number of theoretical studies about single fiber pull-out tests. They

indicated that the process of interfacial debonding can be divided into four stages:

1) linear load-displacement relationship up to the initiation of debonding, 2) stable

cracks propagation with friction, 3) cracks propagation arrives at a limit and the

whole interphase is debonded, 4) frictional sliding along fiber surface (Fig. I-17(a)).

The debonding force should be the force at which micro-cracks were initiated (Fd)

but not the maximum force (Fmax). However, the works of other researchers do not

show the second stage and the maximum force (or peak force) is considered as

debonding force (Fd) (Fig. I-17(b)) [Gaur1989, Biro1991, Craven2000, Zinck2001,

Kang2009, Liu2011]. The different toughness of the interphase layer can explain

the different behaviors. An interphase having a higher ability to be deformed tends

to display yielding to debonding with crack propagation. On the opposite, for a stiff

interphase, the crack propagation achieves instantaneously leading to fast fracture.

The signature of the force-displacement curve is significant of the interphase.

Moreover, different interfacial adhesion strengths can also lead to different

failure modes. Yang et al. [Yang2010] studied the interfacial strength of a glass

fiber/polypropylene system by micro-droplet debonding test and found that the

results can be divided into two categories according to whether or not there was

stable interfacial friction after debonding (Fig. I-18). The observation with the

optical microscopy indicated that, for the samples with a constant friction after

debonding, there is little residual matrix around the debonded area of the fiber (Fig.


38

I-18(a)). This failure mode is denoted as adhesive interfacial failure (category A).

However, for another category (category B), for which the friction decreases

gradually (Fig. I-18(b)) and there is more residual matrix, it is denoted as cohesive

matrix failure. The cohesive matrix failure is considered as a stronger interfacial

strength [Biro1991, Yang2010].

(a) (b)

Figure I-18. Two typical load-displacement traces of micro-droplet debonding test: (a) with a stable

interfacial friction and (b) with a gradually decreasing interfacial friction [Yang2010].

B.2.1.4 Interfacial failure criteria

The primary objective of the single fiber pull-out test is to determine intrinsic

parameters which can characterize the interfacial mechanical properties or, in

other words, that can be used as an interfacial failure criterion. These parameters

should not be affected by the inherent properties of the fiber and/or the matrix, the

sample size or shape, and should be independent on the testing techniques

[Zhandarov2006]. This section will introduce four specific interfacial mechanical

parameters, including apparent interfacial shear strength (τapp), local interfacial

shear strength (τd), critical energy release rate (Gic), and adhesion pressure (σd).

Apparent IFSS (τapp)

Apparent interfacial shear strength is the most widely used parameter to

characterize the interfacial mechanical properties, which can be calculated as

following:


39

𝜏𝑎𝑎𝑎 = 𝐹𝑚𝑚𝑚𝑆

= 𝐹𝑚𝑚𝑚2𝜋𝑟𝑓𝑙𝑒

(2),

where Fmax is measured maximum force, S is embedded area, rf is fiber radius, and

le is the embedded length. Zhandarov [Zhandarov2006] indicated that this

parameter is relevant to qualitatively estimate the interfacial bonding strength or to

compare the interfacial strength of different interphases. However, it can’t be

regarded as an interfacial failure criterion for two reasons: 1) many studies have

reported that τapp evaluated from this simple analysis is not a constant but

decreases with the increase of le [Gaur1989, Scheer1992, Day1998, Kim1998]; 2)

during debonding, interfacial friction contributes to the calculated IFSS value

[Zhandarov2000, 2000a].

Local IFSS (τd)

Local interfacial shear strength used as an interfacial failure criterion supposes

that the interfacial failure occurs when the local interfacial shear strength reaches a

critical value τd [Zhandarov2002, 2005, 2005a]. The current shear force (F) during

crack initiation and propagation could be expressed as following:

𝐹 = 2𝜋𝑟𝑓𝛽

�

𝜏𝑑𝑡𝑡𝑡ℎ[𝛽(𝑙𝑒 − 𝑡)] −𝜏𝑇𝑡𝑡𝑡ℎ[𝛽(𝑙𝑒 − 𝑡)] ∗ 𝑡𝑡𝑡ℎ �[𝛽(𝑙𝑒−𝑎)]

2�

+𝛽𝑡𝜏𝑓

� (3),

where β is shear-lag parameter, τT is a stress due to thermal shrinkage, τf is

frictional stress and a is current crack length. The exact expression of β, τT and τf

will be presented in Annex I. By assuming a equals zero, the debonding force

which initiates cracks can be estimated.


40

Critical energy release rate (Gic)

Scheer and Nairn [Scheer1992, Scheer1995] indicated that, for a single fiber

pull-out test, the interfacial cracks are initiated and propagated at a constant

energy release rate which is called as ‘critical energy release rate’, Gic. When the

embedded length tends towards infinity, the value of Gic is given as following:

𝐺𝑖𝑖 = 𝑟𝑓2�𝐶33𝑠𝜎𝑑2 + 2𝐷3𝑠𝜎𝑑∆𝑇 + 𝐷32

𝐶33∆𝑇2� (4),

where C33, C33s, D3, D3s are constants which correspond to the properties of fiber

and matrix and specimen geometry. The calculation of these values will be

interpreted in detail in Annex I. rf is the radius of fiber, σd is the interfacial shear

strength and ΔT is the temperature difference between the test temperature and

the stress free temperature. Liu and Nairn [Liu1999, Nairn2000] considered that,

during the debonding test, friction and residual stress are also important effects

and they have revised the function (4) for more accurate results.

Adhesion pressure (σd)

Scheer et al. [Scheer1995] investigated the stress distribution in a micro-droplet

specimen during the debonding test. They indicated that the debonding force can

be divided into two directions, radial stress (σrr) and shear stress (τrz), and the

variation of these two stresses is presented in Figure I-19. It’s obvious that, at the

point where fiber enters the matrix, the shear stress is zero and only the radial

stress contributes to the initiation of interfacial cracks. Zhandarov and coworkers

[Zhandarov2006, Pisanova2001] considered that, at this point, there is a critical

value for radial stress. When the current stress exceeds this value, the


41

micro-cracks are initiated. This critical value was called critical adhesion pressure

(σd) and can be regarded as an interfacial failure criterion. The analytical

expression of this value which is based on shear-lag model is given by Scheer

[Scheer1995] as following:

𝜎𝑑 = 𝐹𝑑𝜋𝑟𝑓

2 = �−𝐷3𝑠∆𝑇𝐶33𝑠

+ �2𝐺𝑖𝑖𝑟𝑓𝐶33𝑠

� (5),

where all the parameters are the same as that in function (4) (Annex I). Zhandarov

et al. [Zhandarov2006] developed this function to a new form, in which frictional

stress was considered.

Figure I-19. Interfacial stress distribution along the embedded length: σrr, radial stress; τrz, shear

stress [Scheer2005].

Zhandarov et al. [Zhandarov2002, 2005, 2005a, 2006] have compared the

interfacial failure criteria mentioned above and indicated that, although these

parameters are accordant with experimental data in some conditions, they are

sensitive, to a certain extent, to specimen geometry, to loading conditions, and to

failure mode, so that any of them could be regarded as an absolute failure criterion.

For example, adhesion pressure (σd) is only valid as a failure criterion for an


42

extremely short crack length (initiation of crack) where Mode I failure (tensile) is

predominant. When the cracks propagate in interphase, the interfacial failure mode

is mainly Mode II (shear). In this case, critical energy release rate (Gic) should be

the best failure criterion. Additionally, for the specimens with a very short

embedded length, local interfacial shear strength (τd) seems to yield the best fit

with experimental data. For these reasons, choosing an appropriate parameter to

evaluate interfacial performance appears to be very important. Zhandarov

[Zhandarov2006] considered that the generalized work of adhesion could be

regarded as a universal failure criterion which is applicable under any conditions.

Unfortunately, there is no method which can measure this value accurately.

Figure I-20. Schematic representation of the single fiber fragmentation test [Herrera-Franco1992].

B.2.2 Single fiber fragmentation test

The single fiber fragmentation test was early used to characterize the mechanical

properties of metallic composite material [Kelly1965], which refers to embedding a

single fiber with a finite length in the middle of a dog-bone shaped matrix (no


43

matter metal or polymer). As an external uniaxial stress is applied, a tensile stress

is transferred to the fiber via interfacial shear. When the tensile stress on fiber

increases up to the failure stress of the fiber, the fiber breaks. Continued tension

will result in that the length of fiber fragments become shorter and shorter until the

whole specimen breaks (Fig. I-20). The final fragmentation length of the fiber is

regarded as a critical length which can be used to evaluate interfacial mechanical

properties. To perform the experiment, the strain at break of the matrix must be at

least three times higher than that of fiber [Herrera-Franco1992, Drzal2002].

Generally, the embedded fiber breaks into dozens of sections with various lengths.

Therefore statistical methods are required to evaluate the critical fragmentation

length. For the composite system formed with a transparent matrix, the length of

fragmentation can be measured directly by optical microscopy. However, for some

opaque materials, the optical method seems to be useless. Acoustic emission

technique (AE) is one of the recently developed non-destructive examination

methods, which can be used for the measurement of fiber fragment length

distribution [Netravali1989]. This method is very sensitive to the micro cracks in

material and, however, easy to be influenced by geometrical parameters of sample,

dominant frequency of the emitted ultrasonic signal and, more important, the

deformation of the material.

B.2.3 Single fiber push-out test or micro-indentation test

Single fiber push-out test (or micro-indentation test) was proposed by Mandell

[Mandell1980] in 1980. This test involves applying a compressive force to push out

an extremely short fiber section from matrix block by using a hard micro-indenter.

The variation of compressive force as function of indenter displacement distance is


44

registered, which can be used to evaluate interfacial mechanical performance (Fig.

I-21).

Figure I-21. Schematic diagram of micro-indentation test [Tanoglua2000].

This test is conducted on real composite material and therefore reflects the real

conditions as it takes into account the presence of surrounding fibers. However,

this test has not been widely applied because of the difficulties of specimen

preparation and data analysis. For successfully carrying out this test, the

transverse section of the composite (perpendicular to the fiber direction) should be

subtly polished and only the fibers which are aligned and perpendicular to the

section can be chosen for testing. In addition, the rigidity of the fiber also

influences the testing result. For example, a fiber with a weak hardness will be

indented by the indenter. This part of indentation distance can’t be regarded as the

deformation of interphase, which leads to some mistakes to the testing results. For

this reason, choosing an appropriate indenter is likely to be very important for the

test (Fig. I-22) [Wu1997, Tanoglua2000]. Moreover, the investigation of Tanoglua

and Foley [Tanoglua2000, Foley2002] indicated that the interfacial strength and

energy-absorbing capability was sensitive to displacement rate of indenter. A

higher displacement rate may lead to an over-evaluation of interfacial

performance.


45

Figure I-22. Schematic drawing of indentation test: using a spherical indenter (a) and a Vicker

micro-hardness indenter (b) [Kim1998].

B.2.4 Compression test

Compression test is one of the methods developed in early stage to measure the

interfacial bond strength of composite material, which involves applying a

compression at two ends of specimen, in which a single fiber was embedded along

the direction of loading. By changing the type of specimen geometry, different

interfacial failure modes could be obtained. Figure I-23 displays two different

specimen geometries for compression test. When the specimen, as shown in

Figure I-23(a), is loaded in longitudinal compression, shear stress is generated

because of difference of elastic properties between fiber and matrix. However, with

the second geometry (Fig. I-23(b)), a tensile interfacial debonding occurs under

longitudinal compression, which is mainly caused by the expansion of the matrix,

i.e. the Poisson ratio of matrix is generally greater than that of fiber. In this test,

accurately aligning the fiber is very difficult and the interfacial failure should be

detected visually. For this reason, this method has not been as popular as other

micromechanical methods.


46

(a) (b)

Figure I-23. Single fiber compression test with (a) parallel-side and (b) curved-neck specimen

[Kim1998].

B.3 Conclusion

In the previous sections, almost all the micron-scale testing methods involved in

characterizing interfacial mechanical properties were introduced. However, this

review is not exhaustive but detailed the relevant methods for the following study

dedicated to the interfacial self-healing. For our purpose, i.e. demonstration of the

interfacial healing ability of the designed interphase, the micromechanical methods

should consider the following specificities:

1- Testing target should be only (or as much as possible) related to the

interphase. In this thesis, we just focus on modifying interphase to get a

healing ability. According to this objective, other fracture modes, such as fiber

breakage and matrix fracture, should be avoided during test.

2- The overall structure/shape of the specimen should not be changed after the

first mechanical test. In fact, for characterizing the interfacial healing ability,

the interfacial mechanical properties before and after healing, generally,

should be compared. However, some microscopic methods such as single

fiber fragmentation test which requires to stretch sample until entire fracture.

As a consequence, this type of test will not be selected. Although some


47

methods such as single fiber push-out test will not seriously change the

sample structure and for which interfacial self-healing is easy to be activated,

it does not allow to carry out a second test to evaluate self-healing. Because,

after the test, one end of the testing fiber generally retracts into matrix. The

measurement of fiber embedded length will become impossible for the

second indentation. Besides this, the indenter is also inadequate because of

the tip diameter.

3- Sample should be easy to be prepared. In most cases, the interfacial tests

need that fiber orientation in matrix should be controlled. However, it’s not

easy to be achieved, especially for the micron-scale tests, i.e. the orientation

of some very soft and thin single fibers will be extremely difficult to be

controlled in a viscous matrix.

Among all the methods introduced above, it’s only single fiber pull-out test that

possesses these characteristics. During the single fiber pull-out test, the

micro-cracks propagate principally along interphase, which can exactly display the

interfacial mechanical performance. After the first test, matrix droplet will not

separate from fiber surface but slide along it, which is favorable for interfacial

self-healing. In addition, there is no need to specially control the shape of matrix

droplet which will be formed automatically under the influence of gravity and

wettability.

C. Self-healable composite materials

In the last two sections, we have focused on physicochemical routes for

modification of interphase and on the description of micromechanics experimental


48

approaches which are relevant for evaluating the impact of the modified interphase

on the interfacial adhesion. In this section, we are going to focus on the

self-healing strategies which are reported in the literatures and could be applied to

carbon fiber/epoxy interphases. How to introduce the self-healing ability to the

interphase of a composite to make responsive interphase will be the primary focus.

So far, there are two strategies which have been applied to design self-healing

composite materials, i.e. extrinsic self-healing and intrinsic self-healing

[Zhang2012a]. The extrinsic self-healing consists in embedding some liquid

healing agent within the matrix. As micro-cracks cross this liquid area, the healing

agent will be released and fill the crack. The intrinsic self-healing is achieved from

physical and/or chemical reversible interactions. This part will detail the two

strategies in order to select the best route which will be the most relevant for

self-healable interphases.

C.1 Extrinsic self-healing for polymer materials

The extrinsic self-healing is a bio-inspired strategy (simulation of the bleeding

process) that involves the embedding of some healing-agent-filled hollow

structures such as micro-capsules, hollow fibers as well as three dimensional

microvascular networks into the polymer matrix. When a micro-crack grows and

crosses the filled hollow structures, the liquid healing agent will be released and fill

the crack and then homopolymerizing or reacting with reactive species

(comonomer, grafted functions, etc.). After the polymerization or copolymerization

step, the cracks will be repaired or will be stopped.


49

C.1.1 Embedded microcapsules

Microcapsule embedding method was firstly proposed by White et al. [White2001].

This route involves embedding encapsulated liquid healing agent into polymer

matrix. When there are micro-cracks crossing these microcapsules, the healing

agent will be released under the capillary effect (Fig. I-24).

Figure I-24. Self-healing strategy based on embedding microcapsules into the polymer which

contains the corresponding polymerization catalyst (Grubbs’ catalyst) [White2001].

This technique requires firstly a suitable shell material which should be neither too

rigid nor too fragile. If the capsule shell is too rigid, it will not be fractured and will

not release the healing agent in time during the propagation of cracks. Inversely, if

it is too fragile, the most of the capsules will be broken during processing. For this

reason, urea-formaldehyde system was chosen to encapsulate liquid healing

agent by in-situ emulsion polymerization. The diameter of the capsule is about 50


50

to 200 µm. Besides this, the healing agent is also very important. White et al.

[White2001] have proposed dicyclopentadiene (DCPD) as healing agent and

Grubbs’s catalyst as hardener which is mixed into the matrix prior its processing

(Fig. I-24). This healing agent system presents a long shelf life, a low viscosity and

volatility, fast polymerization at room conditions and a low shrinkage upon

polymerization. Furthermore, the activity and stability of the Grubbs’s catalyst will

not be changed when mixing with epoxy matrix. With this healing agent system, an

average healing efficiency of 60% was obtained. Besides providing an efficient

mechanism for self-healing, the existing of DCPD-filled polymeric microcapsules

and Grubbs’s catalyst particles in matrix can significantly improve the virgin

mechanical properties of the composite [Brown2002, 2005].

Brown and White [Brown2002, 2005] have also studied the influence of

capsule size on healing efficiency. The results confirmed that the bigger the

capsule is, the more DCPD monomer is released into the cracks plane. However,

the variation of the quantity of DCPD monomer in cracks can’t lead to an obvious

change of healing efficiency. In addition, they have also studied the healing ability

of this system under dynamic loading (fatigue test) and concluded that a significant

crack arrest could be resulted when the in-situ healing rate was faster than the

crack growth rate. For satisfying this requirement, a carefully timed rest period

(unloaded) should be applied during loading to ensure a complete polymerization.

Kessler et al. [Kessler2003] also indicated that the healing efficiency of this system

can further be improved by heating during healing process.

Although this method can be used to construct a self-healable polymeric

composite material, there are some unavoidable limitations such as: 1)

constrained by the itself encapsulation technique, since only few of the polymeric


51

monomers can be used as healing agent which is, generally, not the best choice

for adhesive material; 2) this self-healing method can’t work repeatedly in a given

location of the composite [Toohey2007]. After a healing process, the liquid healing

agent near the micro-cracks is almost released completely and can’t be active

again when a crack propagates again; 3) after healing, some voids will be formed

because of liquid releasing, which will lead to defects in composite and will weaken

its mechanical properties [Trask2007, Zako1999]. For these reasons, this method

has not been used widely in the composite materials.

C.1.2 Embedded hollow fibers

The self-healing concept based on embedding hollow fibers was firstly introduced

into composites in 1996 by Dry [Dry1996]. This method is based on a similar

principle as micro-capsule embedding, i.e. micro-cracks trigger the releasing of

liquid healing agent who has been embedded in matrix and then repair the cracks

or arrest their propagation under a simple reaction (Fig. I-25).

Figure I-25. Schematics of the self-healing concept based on hollow fiber embedding [Bleay2001].

Nevertheless, this method has some other unique advantages. First of all, Motuku


52

et al. [Motuku1999] investigated different tube-like materials for potential use as

storage material such as borosilicate glass micro-capillary pipets, flint glass

Pasteur pipets, copper tube, and aluminium tube. The results demonstrated that

glass tube might be a better candidate for this kind of self-healing concept. The

glass tube processing technology has been already very experienced and the

diameter (external and internal) can be controlled according to the composite

material uses. The characteristics of the glass tube allow more possibilities for

self-healing. For example, filling the hollow glass fiber with liquid healing agent is

relatively easy to be achieved (under vacuum) offering a wide choice of monomers

[Trask2007]. Moreover, the glass fiber is a conventional reinforcing material in

composite material. An appropriate embedding of some healing agent within the

hollow glass tube can supply not only a self-healing ability but also a reinforcement

of the composite [Bleay2001, Pang2005a, Trask2006, 2007]. By comparing with

the conventional glass fiber, the external diameter of hollow glass fiber is relatively

larger. Its reinforcement effect in a composite material will be very limited

[Trask2007a]. Additionally, this method will be also non-valid when cracks

reappear in the same location where the healing agent has been already

consumed. Pang and Bond [Pang2005] underlined that this method can’t be

proposed as a permanent self-healing capacity to completely repair the cracks in a

composite but to arrest further propagation.

C.1.3 Tri-dimensional microvascular network

Although the self-healing material composed of micro-capsule or hollow glass fiber

exhibits an excellent healing ability, this function is available only one time in a

given location because of the consumption of healing agent. Furthermore, the


53

outflow of healing agent will create some voids inside the material which will be

also a potential risk for mechanical properties. For these reasons, Toohey and

co-workers [Toohey2007, 2009] have built a new healing system which was based

on the embedding of a tri-dimensional microvascular network instead of the

micro-capsules or hollow glass fibers (Fig. I-26).

Figure I-26. Self-healing polymer-based material including a 3D microvascular circulation network.

a) schematic of a cross-cut on human epidermis layer (bleeding) after self-healing; b) schematic of

a self-healing polymer-based material which contains a biomimetic ‘blood circulation system’

[Toohey2007].

This microvascular is similar to human blood circulation system, by which the liquid

healing agent can be continuously transported to any fractured location and then

fill the cracks. With this healing system, a crack can be healed more than 7 times

with a relatively high healing efficiency. Since the presence of microvascular

impacts the structural properties of the matrix, a network with maximum channel

spacing and minimum channel diameter is required. For a given healing cycle, the

amount of catalyst in matrix has no significant influence on healing efficiency.


54

However, the concentration of catalyst has a clear impact on the number of

successful healing cycles. With a consecutive healing cycle, the available catalyst

amount becomes less and less, and, at the same time, the healing efficiency

decreases gradually until all the catalyst is depleted. At this time, the healing ability

is unavailable, even though the healing agent is supplied continuously.

Figure I-27. Development of microvascular self-healing network [Blaiszik2010].(1) one-dimensional

(1D) healing system based on hollow fibers filled with healing agent; (2) two-dimensional (2D)

microvascular network healing system consists of two isolated microvascular [Toohey2009a]; (3)

three-dimensional (3D) interpenetrated microvascular network healing system [Hansen2009].

For overcoming this limitation, Toohey et al. [Toohey2009a] introduced another

healing system which is composed with two isolated microvascular (Fig. I-27). One

is filled with epoxy resin and another is filled with liquid amine curing agent. This

healing system can continuously supply epoxy resin and curing agent into a crack

and thus increase the number of healing cycles. The experimental results

demonstrated that 60% of healing efficiency was obtained even after 16 healing

cycles. However, Hansen et al. [Hansen2009] indicated that the healing agent


55

(epoxy resin and curing agent) must migrate and diffusively mix over long distance

within the crack, which will seriously influence the healing efficiency. As a

consequence, a 3D interpenetrated microvascular network was constructed as the

tunnel of healing agent (Fig. I-27). The research of Hamilton [Hamilton2010]

confirmed that this interpenetrated microvascular healing system can be also

applied in a bulk structural material. For reducing the curing time of healing agent,

Hansen et al. [Hansen2011] proposed to introduce a third microvascular network in

which a temperature-controlled fluid is introduced so that the damaged region can

be heated locally and thus accelerating the self-healing.

Although the microvascular healing method is developed all the while, there

are some limitations such as: 1) after each healing cycle, a solidified epoxy film

formed by the reaction of healing agent covers the micro-channels, which prevents

the penetration of healing agent during the next healing cycle. That is the reason

why the healing efficiency decreases as the number of healing cycles increasing; 2)

since the penetration or flow of the epoxy resin and curing agent is random, it is

impossible to mix uniformly with an appropriate stoichiometric ratio, which reduces

the healing efficiency of this healing system; 3) this healing system has not been

introduced into a fiber reinforced composite material because of its complex

processing and the influence on final material structure and mechanical behaviour.

C.2 Intrinsic self-healing for polymer materials

Intrinsic self-healing is a relatively passive healing method. It can’t be activated by

the formation of cracks but needs some additional stimulation such as heating or

pressure. Thus, it can’t be considered as an intelligent self-healing to some extent.

However, since this healing principle mainly depends on intermolecular


56

interactions (physical or chemical) such as the rheology of thermoplastic polymer,

reversible covalent bonds, hydrogen bonds, and metal ligand, etc., it can be

realized in many systems by different processing technologies and chemical

modifications. Therefore, this kind of healing method has been widely studied.

C.2.1 Self-healing based on thermoplastic polymer

It’s well-known that the thermoplastic by itself is one kind of self-healable material.

When it’s fractured, a heating above its glass transition temperature (Tg) or a

swelling in appropriate solvents allows macromolecular chains near the crack lips

to interpenetrate and inter-diffuse, i.e. forming ‘chain bridges’ across the cracks

and thus repairing them [Lin1990, Kawagoe1997, Hsieh2001]. However, for a fully

cured thermosetting resin such as epoxy resin, the macro-molecular chains can’t

move similarly as that of thermoplastic polymer, even above its glass transition

temperature. The molecular mobility is strongly limited by the existence of

crosslinks. As cracks appear in such materials, the self-healing concept used for

thermoplastics can’t be adopted. However, the interpenetration and inter-diffusion

capacity of thermoplastic polymer chains is still effective in thermosetting network

and can form some chain bridges between two sides of cracks. For this reason, the

healing ability of thermoplastic can be transferred into thermosetting by blending.

For example, Hayes et al. [Hayes2007] have built a self-healing system based on

a miscible thermoplastic/thermoset blend. Thanks to the excellent compatibility

between the thermoplastic, poly(bisphenol-A-co-epichlorohydrin), and epoxy, even

after curing, the thermoplastic remained dissolved in epoxy network and a healing

efficiency up to 70% was obtained. Their following work also confirmed that this

matrix system can be used to construct a self-healable fiber-reinforced composite


57

material [Hayes2007a]. Nevertheless, the dissolved thermoplastic chains

contribute to decrease the Tg of the matrix material, i.e. decrease the

(thermo)mechanical behavior. Furthermore, the good miscibility between

thermoplastic and thermosetting is a requirement so that the self-healing is

effective. But most of thermoplastic/thermoset blends display a phase separation.

Luo et al. [Luo2009] studied the self-healing ability of an incompatible

epoxy/poly(ε-caprolactone) (PCL) blend. After blending, this system separated into

two phases immediately and formed a ‘bricks and motar’ structure. For keeping

relatively high mechanical properties of this blend, a small quantity of PCL was

added, and thus became a dispersed phase (‘bricks’). Different from the works of

Hayes [Hayes2007, 2007a], the bridging material of this healing system was not

molecular chains but crystallized PCL layers, which contributed to a higher healing

efficiency. When a moderate force was applied to assist crack closure, the healing

efficiency even exceeded 100%.

C.2.2 Self-healing process based on reversible bonds

For the most of polymer materials, the fracture mechanisms proceed from the

breakage of macromolecular chains which is generally irreversible. However, if the

broken macromolecular chains can be reconnected by some reversible bonds or

interactions, the cracks will also be healed. Based on this consideration, many

reversible bonds can be considered such as reversible covalent bonds,

supramolecular assemblies and metal ligand bonds.

C.2.2.1 Reversible covalent bonds

The covalent bond is one considered in conventional polymers (Staudinger’s

macromolecular chain). Therefore, reversible covalent bonds will be the best


58

choice for constructing self-healable polymer materials. Among all the reversible

covalent bonds, the Diels-Alder (D-A) thermally reversible covalent bond is one of

the most widely used nowadays to build thermally self-healable polymers as it can

be separated into a conjugated diene and a substituted alkene above a defined

temperature (retro D-A reaction) and reconnected below this temperature (D-A

reaction). During this process, no additional assistant agent such as solvent or

catalyst is needed. Furthermore, the D-A bond is one kind of covalent bonds which

doesn’t impact the mechanical properties of the polymers formed by them

[Rickborn2004].

Figure I-28. Chemical formulas of 1,8-Bis(maleimido)-3,6-dioxaoctane (2ME), 1,

8-Bis(maleimido)-1-ethylpropane (2MEP), furan monomer (4F), and the reversible D-A reaction

[Chen2002].

The first self-healable polymer based on D-A reversible covalent bonds were

synthesized by Chen et al. [Chen2002, 2003] and were denoted as 2MEP4F and

2ME4F. All the monomer linkages or cross-links of this kind of polymer are formed

by D-A cyclo-addition between furan and maleimide groups (Fig. I-28). A qualitative

study indicated that cracks in these D-A based materials can be healed effectively

with a simple thermal healing procedure and this healing process can be repeated

several times. Thomas [Thomas2007] systematically and quantitatively studied the

healing ability of the 2MEP4F polymer by double cleavage drilled compression

(DCDC). The experiment results indicated that, after repeated fracture-healing


59

cycles, the entire fracture toughness can be recovered by a simple heating at

temperatures from 85 to 95°C. The author attributed the excellent healing

performance to the fact that D-A covalent bonds are weaker than other bonds in

the polymer backbone, so that, when the sample is stressed excessively, it will

break preferentially along these bonds as compared with the other covalent bonds

in the polymer chains. After heating, these broken D-A bonds will be completely

reconnected and thus a macromolecular structure almost similar to the original

polymer will be restored leading to recover the overall macroscopic fracture

resistance. The broken D-A bonds could, in theory, reconnect themselves at any

temperature below the retro-D-A temperature (110 °C). In other words, it should be

possible that these cracks might repair themselves automatically at room

temperature. In fact, after fracture, the free macromolecular chains are restricted

near the planes of cracks, which makes difficult to meet each other as an intimate

contact is required. Thus, an appropriate heating is required to provide sufficient

conformational motion to free macromolecular chains. In addition, a reasonable

pressure is also necessary to accelerate the reconnection of D-A bonds from the

(re)establishment of an intimate contact between crack lips. A similar healing

system was also proposed by Yoshie et al. [Yoshie2010] by combining

furyl-telechelic poly(ethylene adipate) (PEAF2) and tri-maleimide (M3). Since this

system utilizes long-chain telechelic prepolymer (PEAF2: Mn=8.7X103) as

monomers, the concentration of D-A bonds in matrix remains too low. For this

reason, the healing process requires more time to reform D-A bonds and the

healing efficiency is not remarkable.

The self-healing concept reported previously based on Dies-Alder thermally

reversible reaction can be also applied to some conventional thermosets such as


60

epoxy. Tian et al. [Tian2009] synthesized an epoxy monomer which contains both

furan and epoxy groups, N, N-diglycidyl-furfurylamine (DGFA). An epoxy network

was designed from the reaction of DGFA with N, N’-(4, 4’-diphenylamine)

bismaleimide (DPMBMI) and methylhexahydrophthalic (MHHPA) (Fig. I-29). The

resulting epoxy network contains not only thermally reversible Diels-Alder covalent

bonds but also thermally stable bonds that combines both the mechanical

performance of conventional epoxy and also the ability of thermal self-healing at

the same time.

Figure I-29. Chemical formula of N,N-diglycidyl-furfurylamine (DGFA), N,N’-(4,4’-diphenylamine)

bismaleimide (DPMBMI) and methylhexahydrophthalic (MHHPA) [Tian2009].

Peterson et al. [Peterson2010] reported another D-A bond based healable epoxy

network which involved a functionalized epoxy-amine thermoset with furfuryl

glycidyl ether (FGE) to graft a great number of furan groups on epoxy network.

Cracks can be healed by a bismaleimide/DMF solution at room temperature under

minimal pressure. The healing process can be explained as two steps: 1) the BMI

solution penetrates into cracks and swells the matrix, leading to increase the

mobility of furan-functionalized chains; 2) the two bismaleimide functionalized ends

react with furan groups at each side of crack and leading to heal the cracks.

According to their experiments, they also concluded that the loading recovery

contributes by both physical (28.4%) and chemical bonding (41.6%). The physical

bonding is caused by solvent-mediated swelling and interlocking, whereas the


61

chemical bonding results from the formation of thermally reversible D-A bonds.

Although about 70 % of healing efficiency in average is obtained by this method, it

is not a real self-healable epoxy system but it can be considered as an inspiration

for self-healable epoxy.

Based on the Diels-Alder and retro-Diels-Alder reactions, all the neat

polymers already mentioned possess more or less a healing ability. Several works

have focused on the building of a self-healable composite material [Wang2007a,

Park2008, 2009, 2010]. For example, Park et al. [Park2010] made a thermally

self-healing composite by combining the resin composition, 2MEP4F, with carbon

fabric. They obviously observed that the cracks produced by flexural test could be

merged visually in 15 min of heating at 100 °C. However, the mechanical strength

did not achieve the desired result, i.e. the strength after healing was much lower

than that of the virgin material. They attributed this result to the interfacial

debonding and fiber breakage which are irreparable despite the D-A reversible

covalent bonds in matrix backbone.

C.2.2.2 Supramolecular assemblies

Some supramolecular assemblies based on hydrogen bonding and π-π stacking

etc. also exhibit a self-healing potential in polymer materials. The supramolecular

assemblies generally exist under the form of pairs, donor/receptor. These donors

and/or receptors are connected to the polymer chains ends (telechelic polymer) to

construct supramolecular materials which more or less possess a healing ability.

Sijbesma et al. [Sijbesma1997] chemically connected 2-ureido-4-pyrimidone

to chain ends of oligomers, by which a new type of reversible self-assembling

polymer (linear or crosslinked) was constructed under the effect of quadruple

cooperative hydrogen bonds (Fig. I-30). Although this system was assembled by


62

non-covalent bonds, its viscoelastic properties are similar to the conventional

polymers, even in semi-dilute solution. Then, the coworkers of Sijibesma

[Brigitte2000] connected the same hydrogen bonds at chain ends of some

conventional polymers in order to significantly decrease the melt viscosity without

losing their solid-state mechanical properties. For increasing the miscibility

between poly(ether ketone) (PEK) and poly-isobutylene (PIB), Binder et al.

[Binder2004, 2005] also introduced multiple hydrogen bonds between PEK and

PIB to construct a new class of pseudo-block copolymer. In this copolymer, the

macroscopic phase separation could not be observed.

(a) (b) Figure I-30. (a) Quadruple cooperative hydrogen bond formed by 2-ureido-4-pyrimidone and (b)

schematics of self-assembling polymer architecture formed by bi-functional and tri-functional

oligomers [Sijbesma1997].

Although a great number of reversible hydrogen bonds were introduced into the

backbone of polymers, their self-healing ability was not always mentioned. Cordier

et al. [Cordier2008] designed and synthesized one type of crosslinked

self-healable elastomers from vegetable oil fatty acid derivatives,

diethylenetriamine, and urea. All the chains and crosslinks of this elastomer were

associated via hydrogen bonds (Fig. I-31). Comparing with conventional

cross-linked elastomers, this system can’t indefinitely sustain stress without creep

and the strain recovery is slow. However, it possesses a unique ability of

self-healing as it is broken or cut at room temperature. Furthermore, the healing


63

rate and healing efficiency are very considerable and the breaking-healing process

can be repeated many times. Montarnal et al. [Montarnal2008, 2009] also studied

the synthesis methods of this type of self-healable elastomers and described how

to control the molecular architecture.

Figure I-31. Chemically modified fatty acid derivatives for designing self-healable supramolecular

elastomers by two steps: 1) acid groups condensed with a controlled excess of diethylenetriamine;

2) the obtained product reacted with urea [Cordier2008].

Kundu et al. [Kundu2008] developed a supramolecular hydrogel consisting in an

oligomeric electrolyte, poly[pyridinium-1,4-diyl-iminocarbonyl-1,4-phenylene-m

ethylene chloride] (1-Cl). After fracture of the network under high shear stress, the

shear modulus of this hydrogel can be quickly recovered at room temperature,

which was attributed to the chlorine ion that stable hydrogen bonds formed

between oligomer and water (Fig. I-32).

Figure I-32. One of the possible forms for the chlorine-ion mediated H-bonding structure of 1-Cl in

water [Kundu2008].


64

In the conventional polymers, the π-electron-rich and π-electron-deficient groups

can self-assemble with reversible π-π stacking interactions. Greenland

[Greenland2008] and Burattini [Burattini2009, 2009a] developed a supramolecular

self-healable polymer based on π-π stacking interactions (Fig. I-33). This

self-healable polymer involved a blending between a π-electron-rich pyrenyl

(π-electron donor) terminated polyamide or polysiloxane and a polyimide which

contained multiple π-electron-deficient receptor sites along its backbone. When

the π-electron donors and the π-electron receptors encounter with each other, π-π

stacking interactions take place and form a novel material whose mechanical

properties are better than its virgin components showing thermo-responsive

healing ability. The healing ability was attributed to the π-π stacking interactions

that tend to disconnect at elevated temperatures and to randomly reconnect upon

cooling.

Figure I-33. π-π stacking interactions between electron-rich and electron-deficient groups (left) and

supramolecular polymer model based on π-π stacking interactions (right) [Greenland2008].

C.2.2.3 Metal ligands

Metal ligand can be a constitutive element of reversible supramolecular structure

based on a metallic center (metallic atom or ion) and several hetero-aromatic

ligands. Although, so far, there is no self-healable polymer material based on metal

ligand, the potential of this reversible junction in the field of self-healing material


65

has never been ignored. Lohmeijer et al. [Lohmeijer2002] constructed a serial of

block copolymers which were assembled via metal ligands. Their works confirmed

that metal ligands could be effectively used in the polymer design by synthesizing

a polymer within plenty of metal ligands. By using single-molecule force

spectroscopy (SMFS) technique, Kersey et al. [Kersey2007] showed that a metal

ligand assembly separated by mechanical force has the ability of self-assembling.

The investigation of Chow et al. [Chow2007] also indicated that the element (ligand

components and metal center) of metal ligands have the ability of interchanging

during blending. Their conclusions clearly display that metal ligand is a notable

constituent for healing cracks in polymer materials, i.e. junction reforming and

element exchanging may reconnect the broken macromolecular chains at the two

lips of crack. Varghese et al. [Varghese2006] firstly introduced metal ligands into

slightly cross-linked hydrophilic polymeric gel and proved the healing ability of this

kind of material. However, a hydrophilic polymer gel is a polymer with a very weak

mechanical strength which prevents to be used as an engineering material.

Figure I-34. Self-healing principle of polymer materials based on metal ligand [Williams2007].


66

Williams and coworkers [Williams2007, Bielawski2007] developed one kind of

organo-metallic polymers that included of N-heterocyclic carbine (NHCs) and

nickel, palladium or platinum metals. These polymers exhibited not only a high

conductivity but also a crack healing ability at liquid state or in a solid film. The

healing principle is shown in Figure I-34.

C.3 Selection of a strategy for interfacial self-healing

In the first two sections, extrinsic and intrinsic self-healing strategies were briefly

overviewed. Even though some researchers [Bleay2001, White2001, Hayes2007a,

Wang2007a, Park2009, and Peterson2010] tried to combine the healable matrix

with fibers to construct so-called ‘self-healable composite’, the healing ability of

them was not impressive. For example, the average healing efficiency of the D-A

modified epoxy matrix was up to 70%. However, for a composite material

composed with this epoxy matrix, the healing efficiency is only 49.6%

[Peterson2010]. This obvious contrast can be explained from the fracture mode of

composite materials. In general, the fracture of fiber reinforced composite includes

three modes: matrix cracking, fiber breakage, and interfacial debonding. These

fracture modes usually occur at the same time. About the matrix cracking, all the

methods mentioned-above can deal with. Fiber breakage appears usually at the

later stage of crack propagation. Thus, fiber breakage can be avoided if the cracks

can be arrested or healed at their initial stage. About interfacial debonding, simply

combining self-healable matrix with reinforcing fibers will not be effective as there

is no available healing bond or molecular ‘bridge’ could be reformed between fiber

surface and matrix.

The method of embedding healing agent into composite (extrinsic


67

self-healing) seems to have the prospect to solve this problem. Unfortunately, it’s

not the case. For the micro-capsule and micro-hollow fiber embedding method,

both of them require cracks to cross the liquid region. However, the formation and

propagation of micro-cracks in interphase is mainly along the fiber direction, in

which the fracture probability of healing agent capsule or fiber is very low. Maybe

the micro-vascular network system would be more efficient, but the design and

processing remain very difficult.

Obviously, directly combining self-healable matrix with fiber to construct a

self-healable composite is not available, especially for healing the cracks in

interphase, but the healing principle can be referenced. For the extrinsic methods,

the basic principle is filling the cracks with healing agent and curing it so that a

transitional layer adhered to two sides of crack is formed. The same artifice is also

suitable to join the separated fiber surface and matrix again to form a new

interphase.

Figure I-35. SEM micrographs of glass fiber surface after coating with micro-capsules (ξ denotes

the capsule coverage ratio) [Jones2013].


68

Jones and coworkers [Jones2013] are the first researchers who applied this

concept to build an interfacial self-healable composite. Their interfacial self-healing

system involves the coating (sizing) of glass fiber with resin-solvent microcapsules

using a dip-coat technique (Fig. I-35), and then making epoxy micro-droplet on the

coated fiber surface to form an interphase. The interfacial healing ability was

evaluated by the micro-droplet debonding test. After the first debonding test,

almost all the capsules in the region of interphase were broken and the healing

agent (epoxy + solvent) was released between fiber and matrix. Under the swelling

effect of solvent, the residual amine groups of matrix would be moved to react with

the liquid epoxy. After 24 hours of curing process, the specimens were re-tested

and an excellent recovery of interfacial shear strength (IFSS) was achieved.

Moreover, the capsules in interphase did not significantly influence the interfacial

mechanical properties of initial material. The experimental results also

demonstrated that healing efficiency of this system corresponded to the ratio of

capsule coverage and the resin-solvent ratio. The more the capsules covered the

fiber surface, the higher the healing efficiency was.

Since fiber coating techniques can be applied to all kinds of fibers, this

interfacial self-healing method can be introduced in any fiber/matrix composite

systems. Nevertheless, it’s easy to find that the self-healable interphase displayed

here is one-off. If there are secondary damages in material, the healing ability will

not be available because of the lack of healing agent. The healable interphases

based on other extrinsic healing methods such as embedding hollow fibers and

microvascular networks have not appeared. But, predictably, these methods are

also unsatisfactory because of the lack of healing agent after several

breaking/healing cycles and the very complex manufactory and processing.


69

The principle of intrinsic method involves in bridging crack with polymer or polymer

macromolecular chains. These bridges are formed ether by the penetration and

entanglement of polymer chains (self-healing based on thermoplastic polymer) or

by the reformation of interactions between the broken chains between two sides of

crack (self-healing based on reversible bonds). Since the reinforcing fiber surface

is non-penetrable, the self-healing based on inter-diffusion will not be feasible.

However, another route of self-healing based on reversible bond seems to be

convenient. Fiber surface modification introduced in section A.2.2 can be used to

graft some functional groups on fiber surface, by which reversible bonds could be

formed with specific groups on the macromolecular chains of the polymer matrix.

These combined strategies permit to form a self-healable interphase, i.e. to design

a specific self-healing strategy for fiber-based composite materials. This concept

was used by Peterson and coworkers to design a self-healable interphase in glass

fiber/ epoxy composite (Fig. I-36) [Peterson2011].

Figure I-36. Interfacial self-healing based on thermally reversible Diels-Alder reaction

[Peterson2011].


70

From a chemical grafting strategy following a coating process, maleimide groups

were successfully introduced on glass fiber surface. When the modified fiber was

combined with epoxy/FGE/PACM(4,4’-methylene biscyclohexanamine) matrix, the

thermally reversible D-A covalent bonds were formed just in the interphase, i.e.

close to the fiber surface. After interfacial debonding, the IFSS could be recovered

with a simple heating-cooling process. With this system, 41% of average healing

efficiency was achieved and the debonding/healing process could be repeated

more than five times.

Although the healing efficiency of this system is not remarkable, its feasibility

and repeatability is attractive. The possible reasons for relative low healing

efficiency are as following: 1) the concentration of D-A bonds introduced into

interphase is too low; 2) coating process used to grafting functional groups on fiber

surface would be not the best path, because there is no strong adhesion between

coating layer and fiber surface. In this case, interfacial crack is likely to occur

between fiber surface and coating layer. In my opinion, by optimizing the process

of interfacial modification, the potential of this healing system could be developed

further. In fact, the coating was deposited onto a previous sizing which is already

done and required during glass fiber spinning. One can imagine to introduce the

self-healing chemical system into the initial sizing.

Other non-covalent reversible bonds such as hydrogen bonds or metal

ligands can also be introduced into interphase. But their feasibilities should be

proved. For example, multiple hydrogen bonds are generally formed by a pair of

functional groups with very large steric hindrance. When the hydrogen bonds are

separated during interfacial debonding, the mobility of donors and receptors on

fiber surface and in bulk matrix would be an important impact factor for healing


71

efficiency. In addition, if the hydrogen bond pairs are asymmetrically separated, i.e.

other non-hydrogen bonds are broken, are there some ‘switches’ which allow

hydrogen bonds to separate and then randomly reconnect? Using metal ligands to

build a self-healable polymer is still in the primary stage, which takes effect only in

some ‘soft materials’ such as solution or gel. Therefore, it remains difficult to build

a self-healable composite with metal ligands-based approach. According to the

above comparisons and analysis, we decide to choose D-A thermally reversible

covalent bonds as the strategy to build self-healable interphases.

D. Conclusion

In this part, we mainly summarized the background for our research subject,

including: composite interfacial modification, interfacial micromechanics, and

self-healing approaches for polymer materials. From a proper analysis of the

specific requirements associated with the objective of designing self-healing

interphases, we have decided to design self-healable carbon/epoxy interphases

based on thermally reversible Diels-Alder reactions. The maleimide groups will be

grafted on carbon fiber surface by surface modification and the furan groups will be

introduced into epoxy network by using a bi-functional reagent, furfurly glycidyl

ether (FGE). Combining these two basic components together, a carbon/epoxy

interphase based on reversible covalent bonds and having a self-healing ability for

several runs is targeted. For accurately estimating interfacial self-healing ability,

the micro-droplet debonding test is chosen to characterize the mechanical

properties of the interphase. More details about interface formation and

characterization will be given in Part II.


72


73

Part II Design and characterization of a self-healable interphase A. Methodology to generate a self-healable carbon/epoxy interphase ... 75

B. Carbon fiber surface treatments and characterizations ....................... 76 B.1 Introduction ..................................................................................................... 76 B.2 Carbon fiber surface treatment ....................................................................... 78 B.3 Characterization of carbon fiber at different treatment steps .......................... 80

B.3.1 Characterization of carbon fiber after oxidization step ..................................... 81 B.3.1.1 Selection of initial carbon fiber ......................................................................... 81

B.3.1.2 Fiber surface topography after oxidization step ............................................... 83

B.3.1.3 Carbon fiber surface chemical composition after oxidization step ................... 85

B.3.1.4 Influence of HNO3 oxidization on fiber mechanical properties ......................... 87

B.3.2 Characterization of carbon fiber after grafting with TEPA and BMI .................. 89

B.3.2.1 Selection of amination agent ............................................................................ 89

B.3.2.2 Fiber surface topography after grafting with TEPA and BMI ........................... 90

B.3.2.3 Fiber surface chemical composition after grafting with TEPA and BMI ........... 91

B.3.3 Conclusion ...................................................................................................... 94

C. Epoxy matrix composition and characterization ................................... 95 C.1 Composition and curing of matrix system ....................................................... 95 C.2 Matrix characterization ................................................................................... 97

C.2.1 Chemical structure of fully cured furan-functionalized epoxy networks ............ 97

C.2.2 Physical properties of furan functionalized epoxy network .............................. 98

D. Formation and evaluation of thermally reversible self-healable interphase .................................................................................................... 100

D.1 Investigation of Diels-Alder reaction ............................................................. 100 D.2 Formation of interphase containing thermo-reversible bonds ....................... 105 D.3 Interphase micromechanics: micro-droplet debonding test .......................... 108

D.3.1 Specimen preparation ................................................................................... 108

D.3.2 Process of test .............................................................................................. 109

E. Conclusion .............................................................................................. 112


74


75

Part II Design and characterization of a self-healable interphase A. Methodology to generate a self-healable carbon/epoxy interphase

To make self-healable carbon fiber/epoxy interphases, we have chosen to

introduce maleimide groups on the carbon fiber surface which could react with

furan groups included in epoxy matrix. This process includes three steps:

i) activation of the carbon fiber surface to graft maleimide groups,

ii) Diels-Alder bonds formation to design the self-healable interphase,

iii) evaluation of the interfacial healing ability by repeating debonding/healing

cycles (Fig. II-1).

Figure II-1. Strategy of designing a self-healable carbon/epoxy interphase based on D-A reversible

reaction: maleimide group, furan group, D-A bond, epoxy group, amino-hydrogen

group, and irreversible covalent bond.


76

Carbon fiber surface is very inert because of its stable turbostratic graphite

structure. For this reason, the first step which consists in surface functionalization

will be described intensively. In the second step, the content of furan donor,

obviously, influences the structure of epoxy network, and thus directly influences

the mechanical properties of the whole material. The matrix composition and the

interphase formation will be considered carefully. The third step deals with the use

of a micro-droplet debonding test as the micro-mechanical analysis to evaluate the

interfacial mechanical properties and the reversibility of fiber-matrix interface

(interphase), i.e. the self-healing behavior.

B. Carbon fiber surface treatments and characterizations

B.1 Introduction

The ex-PAN carbon fibers used in this work, including T700-99, T700-90, and

T700-50C, were supplied by Toray Inc., and display high tensile strength with

standard modulus. These fibers are proposed to be used in some applications

which require high performance properties of the resulting composite materials

such as pressure vessels, aeronautical parts, etc. Table II-1 displays some basic

properties about T700 carbon fiber [Toray technical data sheet].

Properties of T700 carbon fiber

Tensile strength 4.9 GPa Tensile modulus 230 GPa Strain at break 2.1% Density 1.8 g/cm-3 Filament diameter 7 µm

Table II-1. Properties of T700 carbon fiber [Toray technical data sheet].


77

The main differences between the fibers are that their surfaces have undergone

different treatments. T700-99 is neat T700 carbon fiber, i.e. ex-PAN carbon fiber

after graphitization. Its surface has been neither treated nor protected, and

therefore any pollution should be avoided during manipulation. T700-90 fiber is

T700-99 after oxidization step. Its surface is claimed to contain more active

functional groups than T700-99. Since the functional groups (polar groups such as

carboxyl and hydroxyl) on T700-90 are sensitive to surrounding environment, the

manipulation with this fiber should also be performed carefully. T700-50C is an

epoxy-sized T700-90 fiber. The sizing amount is about 1%0wt. On the one hand,

the coating can effectively protect the active functional groups on fiber surface; on

the other hand, it can improve wettability with reactive epoxy resin and increase

the interfacial strength between fiber and epoxy matrix. For this reason, only

T700-50C is sold as a commercial product and the two other kinds are used just

for investigation (www.toraycfa.com/pdfs/T700SDataSheet.pdf).

According to the results of X-ray photoelectron spectrum (XPS) analysis, we

have found that the surface chemical structures between T700-99 and T700-90

are not very different. Furthermore, we have also tried to graft amino functional

groups on T700-90 directly, i.e. without additional surface treatment. Nevertheless,

the grafting efficiency was not significant. So we have decided to oxidize T900-99

by ourselves to increase the concentration of active groups on fiber surface for

further grafting. After oxidization, the fiber surface will be turned to be

amine-functionalized and grafted with maleimide-terminated compound. The

surface treatment process and analysis of the surface physical/chemical structure

will be described in the following sections.


78

B.2 Carbon fiber surface treatment

The turbostratic graphitic structure of the T700-U * surface is very inert and almost

no functional groups are present for directly grafting. Thus, at first, a surface

oxidization treatment was performed to activate the fiber surface and to generate

polar groups such as carboxylic or carbonyl groups [Gardner1996, Pittmanjr1997,

1997a]. Although T700-U was always stored in a clean plastic bag and

manipulated with rubber gloves and tweezers, as it’s difficult to avoid additional

impurities on its surface. For this reason, a cleaning process was carried out in

refluxed acetone for 1 hour before oxidization treatment.

The oxidization treatment was achieved in a nitric acid solution (HNO3

concentration: 69 vol%) at 115°C** which was proposed as the best conditions for

carbon fiber surface oxidization treatment with nitric acid [Wu1995]. The

oxidization time is considered as a key parameter which not only determines the

activation efficiency but also influences the mechanical properties of carbon fiber.

In fact, the treatment using HNO3 removes in the first step the surface defects

(disordered graphitized layers), i.e. leading to improve mechanical properties,

whereas for excessive oxidization treatment, the mechanical performance of

carbon fibers will be decreased. So, the fibers were oxidized, respectively, for 30,

60, and 90 minutes. Furthermore, the contact of carbon filaments may influence

the oxidization efficiency. For avoiding this case, fiber bundles were loosely put

into nitric acid with an appropriate stirring. After oxidization, the fibers were rinsed

with large amount of distilled water until the pH value of effluent liquid was constant.

Then, the rinsed fibers were dried at 80°C for 24 hours under vacuum and are

* For differing from the trade name of Toray, we used T700-U to replace T700-99 (the same below). ** 115°C is very close to the degradation temperature of nitric acid (about 120°C). A slight temperature

fluctuation may lead to a violent degradation (NO2) which is very harmful for man and environment. Therefore,

necessary protective measures should be set up during this manipulation.


79

denoted as T700-HNO3.

In the further treatment step, T700-HNO3 reacted with

tetraethylenepentamine (TEPA, technique grade, Sigma-Aldrich) (Fig. II-2) at

190°C for 17 hours to graft amino groups which are required for the further Michael

addition reaction with maleimide groups [Hopewell2000]. Since TEPA is a

multi-functional organic amine, all the amine groups (primary and secondary amine)

have the possibility to react with acid groups on fiber surface. Therefore, the

configuration of TEPA molecule on fiber surface could be as ‘single end-grafted’ or

as a loop (Fig. II-2). Additionally, we have also tried to use ethylene diamine (EDA)

as amination agent to react with T700-HNO3 in the same conditions. The results of

XPS demonstrated that the nitrogen contained in the fiber surface after reaction

with TEPA is two times higher than that with EDA, i.e. using TEPA can get higher

amination efficiency. After this step, the fibers were rinsed with large amount of

distilled water until the pH value of effluent liquid was constant. Then, the rinsed

fibers were dried at 80°C for 24 hours under vacuum and are further denoted as

T700-TEPA.

The last step was achieved through the so-called Michael reaction between

amino groups on fiber surface and maleimide groups in

1,1'-(methylenedi-4,1-phenylene)bismaleimide (BMI, purity: 95%, Sigma-Aldrich):

N, N-Dimethylformamide (DMF) solution (Fig. II-2). This reaction is one type of

conjugated additional reactions which involves the opening of C=C bonds on

maleimide group by amine. The reaction condition is at 80 °C for 2 hours. For such

a reaction, the two maleimide groups of BMI molecule have the same reactivity, i.e.

the two groups may react at the same time (back-biting), which is adverse to obtain


80

a maleimide-rich surface. As a consequence, a saturated solution of BMI was used

to avoid back-biting.

After the different treatments, the fibers were washed with DMF for 2 hours to

remove the excess of unreacted BMI that is just physically absorbed on the fiber

surface. The last washing procedure consisted in a rinsing with acetone. In the

final step, the washed fibers were dried at 80°C for 24 hours under vacuum. The

final carbon fibers are denoted as T700-BMI. The chemical structure formulas of

corresponding reaction agents and the different steps of the fiber surface

treatment are summarized in Figure II-2.

Figure II-2. Protocol for grafting maleimide groups on T700-U carbon fiber surface.

B.3 Characterization of carbon fiber at different treatment steps

In order to evaluate the changes introduced by the surface treatment,

characterizations of the neat carbon fiber surface are required. Its topography and


81

roughness were determined by Atomic Force Microscopy (AFM). Its surface

chemical composition was analyzed by Attenuated Total Reflectance Fourier

Transformed Infrared spectroscopy (ATR-FTIR) and X-ray photoelectron

spectrometry (XPS). The tensile strength of treated T700 carbon fiber was

measured by single fiber tensile test. The related characterization methods are

described in Annex II. Besides these methods, we have also tried to use other

methods such as surface element analysis (SEA), nuclear magnetic resonance

(NMR), titration and fiber surface wettability test to characterize fiber surface.

However, the results were not robust enough and thus will not be reported here.

B.3.1 Characterization of carbon fiber after oxidization step

B.3.1.1 Selection of initial carbon fiber

The XPS analysis allows to identify some acidic functional groups required as

potential grafting sites. Table II-2 displays the surface element weight ratio (C, N,

and O) present onto the T700-U surface and the commercially oxidized carbon

fiber (T700-90).

Fiber O% N% C% O/C N/C

T700-U 11.01 1.42 87.57 0.126 0.016

T700-90 10.47 1.68 87.85 0.119 0.019

Table II-2. Fiber surface element percentage determined by XPS.

The surface elementary compositions of the two fibers seem to have no significant

difference. The O/C ratio of the oxidized carbon fiber T700-90 is even surprisingly

lower than that of the raw carbon fiber T700-U. The content of -COOH on fiber

surface was also determined. The peak near binding energy 289 eV is generally


82

attributed to the carbon atom of -COOH group and thus its area percentage can be

considered as the relative content of -COOH on fiber surface [Gardner1996].

Figure II-3 displays the C1s XPS spectra of T700-U and T700-90, respectively. The

contents of -COOH on T700-U and T700-90 are very close to each other.

(a)

(b) Figure II-3. XPS C1s spectra of T700-U (a) and T700-90 (b) carbon fiber surfaces.

These results indicate that the commercial oxidization does not increase the

content of oxygen-containing groups on carbon fiber surface, especially the

content of carboxylic groups. So which carbon fiber should we choose for the

T700-U 288.74eV

3.6%

T700-90 288.99eV

3.4%


83

further grating reaction, commercially oxidized carbon fiber (T700-90) or the fiber

oxidized by ourselves (T700-HNO3). Only the amino treatment allowed to select

the proper carbon fiber. After grafting TEPA amine on carbon fiber surface, the

N/C ratio depended on the content of initial acidic functional groups. The N/C ratio

was only equal to 0.026 for the commercially oxidized carbon fiber (T700-90)

whereas N/C is equal to 0.105 for the raw carbon fiber oxidized for 60 minutes

(T700-HNO3 (60min)). As a consequence, the raw carbon fiber (T700-U) was

chosen to be directly oxidized for further grafting reaction.

Figure II-4. Evolution of fiber surface average roughness measured by AFM as a function of

oxidization time.

B.3.1.2 Fiber surface topography after oxidization step

During the fiber surface oxidization process, HNO3 plays the role not only as an

oxidization agent but also as a strong acid which corrodes fiber surface by

removing the disordered graphite layers. For studying the corrosion effect of HNO3,

surface roughness of the fiber oxidized for different times were measured by AFM.

Figure II-4 displays the fiber surface average roughness as a function of

oxidization time. Below 60 minutes, surface roughness increases linearly with


84

oxidization time. However, the surface roughness decreases for 90 minutes.

These results suggest that before 60 minutes of oxidization treatment, only the

outer surface layer of carbon fiber is etched. While, for 90 minutes of oxidization

treatment, a deeper surface layer was etched creating interconnected sub-surface

pores, internal voids and cracks in the bulk fiber. At the same time, the protrusion

on surface would be smoothened due to the corrosion effect. Figure II-5 displays

the typical AFM images of fiber surfaces after different oxidization time, which

visually shows the change of surface topography attributed to acid etching effect.

Figure II-5. Typical AFM topography of carbon fiber surfaces after different oxidization times.

T700-HNO3(30min) Ra=1.32nm

T700-U Ra=0.20nm




85

B.3.1.3 Carbon fiber surface chemical composition after oxidization step

Since the efficiency of the first treatment step, corresponding to oxidization,

directly determines the efficiency of the two following steps, i.e. yield of grafting,

the influence of oxidization time on the amount of acidic groups generated on fiber

surface must be evaluated. Figure II-6 reports the evolution of ATR-FTIR spectrum

as a function of the oxidization time. After oxidization, the band near 1,700 cm-1

appears, which corresponds to C=O of carboxyl group (-COOH). When oxidization

time is less than 60 minutes, the quantity of -C=O increases gradually. However,

when oxidization time reaches 90 minutes, this amount does not continuously

increase any longer (a slight decrease is observed). Pittman et al. [Pittmanjr1997,

1997a] have observed an increase of the quantity of acidic groups on carbon fiber

surface with oxidization time in presence of HNO3. They also concluded that this

increase was not related to the increase of density of acidic groups per unit area,

but to the increase of fiber surface area arising from corrosion effect of nitric acid.

They observed that, after 60 minutes, the etching rate increased significantly. The

different result obtained after 90 minutes can be possibly attributed to the sample

of measurement, i.e. a bundle of fiber that may lead to an unfaithful result. The

different result obtained after 90 minutes can be possibly attributed to the sample

of measurement, i.e. a bundle of fiber that may lead to an unfaithful result. In

addition, the other functional groups such as -C-H, -C=C-, -CH2-, and -CH<

reducing or undulating (-C-O-C-) easily detectable as increasing of the oxidization

time can be attributed to the corrosion of fiber surface via nitric acid treatment.


86

Figure II-6. Comparison of ATR-FTIR spectra realized on carbon fiber surfaces oxidized for

different times (HNO3, 115°C).

Figure II-7. Relative area of the XPS peaks at binding energy 288.8 eV related to carboxylic groups

as a function of the oxidization time for T700-HNO3 carbon fibers.

For confirming these conclusions obtained by ATR-FTIR, XPS spectroscopy which

is more sensitive was used to characterize the change of the acidic groups on fiber

surface as a function of the treatment time in HNO3. Figure II-7 reports the relative


87

peak area of C1s binding energy (BE) at 288.8 eV, which is related to the presence

of carboxyl groups (-COOH). The amount of -COOH groups increases with

oxidization time even after 60 minutes in agreement with Pittman’s works

[Pittmanjr1997, 199797a].

The XPS results confirm our assumptions, i.e. the content of acidic groups

always increases with oxidization time because of the increasing of surface area;

however, when the oxidization time is higher than 60 minutes, although the total

content increases, the most of them are ‘hidden’ in voids of sub-surface.

Combining the analysis of AFM, the mechanism of carbon fiber surface oxidization

can be schematically displayed as shown in Figure II-8.

Figure II-8. Evolution of the oxidization state of the carbon fiber surface as a function of the

oxidization time.

B.3.1.4 Influence of HNO3 oxidization on fiber mechanical properties

Since the fiber surface is etched by nitric acid during oxidization creating thus

some cavities, the influence of these cavities on the mechanical properties of

single fiber was studied. The study was carried on measuring the fracture tensile

strengths of carbon fiber monofilaments varying with different oxidization times.

The testing method will be detailed in Annex II. For getting accurate results, more


88

than 50 samples of each fiber have been made with two different length, 1cm and

0.5cm. The test results demonstrated that oxidization does not significantly change

the tensile strength at break, at least, within 90 minutes (Fig. II-12).

Figure II-12. Dependence of tensile strength at break of carbon fiber monofilament as a function of

oxidization time.

This result can be attributed to the limited etching depth. According to the results of

AFM, we know that the variation of average surface roughness is only several

nanometers. Compared with fiber diameter, this depth is negligible and thus can’t

significantly influence the mechanical properties of carbon fiber. In addition, we

can also find that the tensile strength at break of short monofilament (0.5 cm) is a

little higher than that of long monofilament (1 cm), which is related to the defect

number in the fiber. The longer the tested monofilament is, the more defects exist

in it and thus the easier will be the fracture [Pardini2002]. This conclusion requires

that the length of fiber between droplet and upper jig (free length) should be short


89

enough to ensure the fiber is not broken before interfacial debonding. The

appropriate length will be discussed in Part III.

B.3.2 Characterization of carbon fiber after grafting with TEPA and BMI

B.3.2.1 Selection of amination agent

Tetraethylenepentamine (TEPA) is a polyfunctional organic amine. In theory,

successfully grafting of one TEPA molecule allows to introduce 6 amino-hydrogen

groups, which is favorable for the further reactions. However, the long chain

molecule can ‘bite back’, i.e. more than one functional group on the same molecule

reacted with fiber surface. In this case, some active acidic groups on fiber surface

can be screened by the TEPA molecules and are not more accessible for grafting,

which can impact the amination reaction efficiency. To confirm this effect, a low

molar mass organic amine, ethylenediamine (EDA), was grafted onto the oxidized

carbon fiber T700-HNO3 under the same reaction conditions. The results of XPS

based on T700-HNO3(90min) are shown in Table II-3. Obviously, utilizing TEPA as

amination agent can get better amination efficiency even if there are back-biting

reactions, i.e. grafting more amino groups on fiber surface (see N/C).

Fiber O% N% C% O/C N/C

T700-HNO3 (90min) 22.53 2.49 74.98 0.3 0.033

T700-HNO3(90min)-TEPA 12.78 7.85 79.37 0.16 0.099

T700-HNO3 (90min)-EDA 13.55 3.83 82.62 0.164 0.046

Table II-3. Elemental compositions of fiber surface after grafting with TEPA and EDA (T700-HNO3

(90min) is the reference).


90

B.3.2.2 Fiber surface topography after grafting with TEPA and BMI

Since both TEPA and BMI are long chain molecules, their grafting on fiber surface

will change surface topography. AFM images reported in Figure II-9 display the

fiber topography after each surface treatment step.

Figure II-9. AFM topographies performed on carbon fiber surfaces after different surface treatments:

oxidization of the T700-U for 60 minutes, amination and BMI grafting (Both TEPA amination and

BMI grafting were performed on T700-HNO3(60min)).

After 60 minutes oxidization, the fiber surface became rougher (convex) due to the

corrosion effect after HNO3 solution exposure. The amination treatment based on

TEPA grafting smooths the surface topography since the long chains of polyamine

cover the fiber surface after grafting. As the TEPA amine functionality is equal to 7,

the molecular architecture of the resulting surface (layer) remains complex.

T700-U Ra=0.21±0.03nm

T700-HNO3 (60min) Ra=2.60±0.64nm

T700- TEPA Ra=2.41±0.45nm

T700-BMI Ra=5.64±3.60nm


91

Nevertheless, the Michael addition reaction between BMI and the primary and

secondary amines leads to an increase of roughness which can be highlighted by

AFM. This phenomenon confirms the grafting of BMI component on carbon fiber

surface as the formation of multi-layers, i.e. TEPA reacts with acidic groups on

fiber surface to form the first layer and BMI reacts with amino groups to form the

second layer.

Figure II-10. ATR-FTIR spectra changes of T700-HNO3 (60min) at different steps of grafting.

B.3.2.3 Fiber surface chemical composition after grafting with TEPA and BMI

The carbon fiber oxidized for 60 minutes can carry more acidic groups on the top

layer which will be favorable for further grafting reactions. Therefore,

T700-HNO3(60min) was chosen as the activated surface for further reaction with

TEPA and BMI in turn. Figure II-10 displays the evolution of ATR-FTIR spectrum

after each grafting reaction. After the amination reaction, the band located at 1,700

cm-1 (characteristics of -COOH) is obviously weakened and some bands such as

1,000 cm-1 (-N< aliphatic) and 3,000~3,250 cm-1 (-NH2 and -NH-) appear. It means


92

that amino functional groups have been grafted on fiber surface successfully. The

BMI grafting reaction will convert most of primary and secondary amine groups into

tertiary amine and, at the same time, will increase the content of -C=O (1,700 cm-1)

on the fiber surface. This analysis is basically in agreement with the results

displayed in Figure II-10 which evidences BMI molecules have been grafted on the

fiber surface.

Fiber

N/C (%)

Relative peak area at 288.8eV (-COOH) (%)

�𝑫𝑵

𝑫−𝑪𝑪𝑪𝑪�

Before amination

After amination

DN Before

amination After

Amination D-COOH

T700-HNO3(30min) 2.9 10.7 7.8 4.9 4.8 -0.1 78

T700-HNO3(60min) 2.2 10.5 8.3 7.0 4.8 -2.2 37

T700-HNO3(90min) 3.3 9.9 6.6 14.5 4.6 -9.9 0.66

Table II-4. Change of nitrogen content and carboxyl group concentrations before and after

amination reaction. DN and D-COOH are the differences of nitrogen content and carboxyl groups

before and after amination, respectively.

The efficiency of TEPA grafting reaction depends on the content of acidic

functional groups generated during the oxidization step. The amination was also

evaluated both by the increasing amount of nitrogen introduced onto the carbon

fiber and by the decreasing amount of carboxyl groups as a function of the

oxidization time. Table II-4 displays the changes of N/C ratio and relative peak

area at 288.8eV (XPS, corresponding to -COOH groups) before and after

amination reaction. Obviously, after amination reaction, the relative content of

-COOH decreases and the content of nitrogen atom increases which means that

TEPA have been successfully grafted on the fiber surface. On the other hand, we

can notice that, for the fibers with different oxidization times, the proportion


93

between the decrease of -COOH and the increase of nitrogen atoms (� 𝑫𝑵𝑫−𝑪𝑪𝑪𝑪

�) does

not correspond with the variation of the quantity of acidic groups on fiber surface.

This effect can be attributed to the ‘back-biting’. For example, when the content of

-COOH on fiber surface is not very high, the distance between two neighboring

-COOH groups is too large for the ’back-biting’ reaction of the TEPA molecule.

Under these conditions, consuming one -COOH group must, in theory, allow to

introduce 5 nitrogen atoms. On the opposite, when the density of -COOH groups

on fiber surface is high, the probability of ‘back-biting’ will increase. In this case,

the number of nitrogen atoms introduced by consuming one -COOH group will

decrease significantly. We have confirmed that the content of -COOH on carbon

fiber surface increases with the increasing of oxidization time and thus the

back-biting becomes more and more significant with the increasing of oxidization

time. This phenomenon can be simply explained by Figure II-11.

Figure II-11. Schematic pictures to illustrate the back-biting reaction happening during the amine

grafting onto the carbon fiber: without ‘back-biting’, consuming 1 -COOH introduces 5 nitrogen

atoms; however, with ‘back-biting’, consuming 1 -COOH introduces only 2.5 nitrogen atoms in

average.


94

Although it seems that, without ‘back-biting’, more amino-hydrogen groups can be

grafted on fiber surface which would be favorable for the further BMI grafting

reaction, the bonding strength of interphase composed with it will be low owing to

the lack of necessary connections between fiber and matrix.

Since the fiber surface became very complex (multiple layers) after BMI

grafting reaction and there is no unique type of functional group on the BMI

molecule, the fiber surface chemical structure can’t be characterized by XPS. More

information about the BMI grafting efficiency will be reflected indirectly by the

interfacial healing efficiency reported in the following Parts.

B.3.3 Conclusion

In this section, we have treated the carbon fiber surface according to three steps,

oxidization, amination, and grafting to get a maleimide-functionalized surface. At

first, the treatment process was described in detail. Then, the physical and

chemical changes of fiber surface composition and topography after each

treatment step were characterized by AFM, ATR-FTIR, and XPS. Furthermore, the

influence of nitric acid oxidization on the tensile strength at break of fiber was also

evaluated by single fiber tensile test. According to the analysis above, some

conclusions are obtained as following:

1) Within 60 minutes of HNO3 treatment, the fiber surface roughness increases

linearly with oxidization time. When fiber is oxidized for 90 minutes, the

protrusion on surface would be smoothened and interconnected pores or

internal voids and cracks can be created in the sub-surface due to the

etching effect.


95

2) The content of acidic group increases with the oxidization time because of

the increasing of surface area and when the oxidization time is more than

60 minutes, most of them are located in the internal surface voids of the

sub-surface. Choosing TEPA as amination agent can get a significant

amination efficiency, able to react with the bismaleimide.

3) Within 90 minutes, oxidization treatment will not damage fiber mechanical

properties as it concerns only the outer surface of the carbon fiber.

C. Epoxy matrix composition and characterization

C.1 Composition and curing of matrix system

Component Chemical formula Product

information

DGEBA H2C

O

HC C

H2O C

CH3

CH3

O CH2

HC

HO

CH2

O C

CH3

CH3

O CH2

HC

OCH2

Dow/ Hunstman 𝑡�=0.15

IPD CH3

H3C

H2N

CH3

NH2

Sigma-Aldrich Purity :99%

FGE O

OO

Sigma-Aldrich

Purity :96%

Curing reaction

CH

O

CH2 NH2CH CH2 N

HOH+

CH CH2 N

HOH

CH

O

CH2 + CH CH2 N

OH 2

k1

k2

Table II-5. Chemical formula of reagents used for epoxy-amine matrix synthesis.

The epoxy-amine system used to design the epoxy matrix was based on a

diglycidyl ether of bisphenol-A epoxy prepolymer (DGEBA) and an isophorone

diamine (IPD) as harderner. In order to introduce furan groups into the epoxy

network architecture, furfuryl glycidyl ether (FGE) monomer containing a furan

group and an epoxy group at each end of molecule was introduced into the


96

reactive DGEBA/IPD system. Thus, the furan-functionalized epoxy matrix was

prepared. The hardener, i.e. IPD, was added with a stoichiometric ratio, amino

hydrogen-to-epoxy equals to 1. The structure of chemicals and corresponding

reactions are shown in Table II-5.

The curing of the DGEBA/FGE/IPD system was performed at 80°C for 2

hours followed by a post-curing at 160°C for 2 hours. Figure II-13 displays a typical

DSC curve of the cured epoxy-amine system. Besides the glass transition, any

other thermal behavior can be observed, which means the matrix system was

completely cured in these conditions.

Figure II-13. Typical DSC curve of a cured epoxy-amine system: DGEBA/IPD (amino-hydrogen :

epoxy =1) (Heating and cooling: +/- 10K/min).

As FGE is a mono-functional monomer for epoxy-amine system (one epoxy group

available for reaction with IPD), it acts as a chain extender between crosslinks, i.e.

FGE contributes to change the crosslink density of the final networks. This change

of epoxy network crosslink density may influence the properties of interphase.

Thus, series of epoxy matrix systems with different contents of FGE were prepared

and shown in Table II-6.


97

System I II III IV V VI

DGEBA/g 10 9 8 7 6 5

FGE/g 0 1 2 3 4 5

IPD/g 2.33 2.35 2.38 2.4 2.42 2.45

Table II-6. Composition of epoxy matrix with different ratios of FGE (amino-hydrogen-to-epoxy =1).

C.2 Matrix characterization

C.2.1 Chemical structure of fully cured furan-functionalized epoxy networks

The furan-functionalized DGEBA-FGE/IPD polymer network was analyzed by

ATR-FTIR. Details about the ATR-FTIR are given in Annex II. Figure II-14 displays

a typical ATR-FTIR spectrum of the cured matrix (system V, DGEBA : FGE = 6:4)

compared with the one from FGE monomer. The presence of the absorption bands

located at 1,503, 1,150, 1,014, and 750 cm-1 corresponding to furan group confirm

that furan groups were introduced in epoxy network by copolymerizing FGE with

DGEBA and IPD.

Figure II-14. Comparison between the ATR-FTIR spectra of furan-functionalized epoxy

(DGEBA/FGE+IPD) network and FGE monomer.


98

C.2.2 Physical properties of furan functionalized epoxy network

As mentioned previously, introducing furan groups within an epoxy-amine network

leads to the increase of molar mass between crosslinks, i.e. decreasing the

crosslink density, which will further influence the mechanical properties of matrix

and also the interphase with functionalized (or non-functionalized) carbon fibers.

For a thermoset polymer, the change of network density can be directly reflected

by the glass transition temperature (Tg). Generally, higher the crosslink density is,

higher the Tg is. Figure II-15 displays the variation of Tg as a function of FGE

content. Obviously, Tg decreases with the increasing of FGE ratio.

Figure II-15. Evolution of the glass transition temperature (Tg) of matrices as a function of mass

ratio of FGE (fully cured DGEBA-FGE / IPD networks).

The dependence of tensile fracture strength (σf) and Young’s modulus (Em) as a

function of FGE content was also evaluated from uniaxial tensile tests. The results

are shown in Figure II-16. When the content of FGE is relatively low (<20%), the

tri-dimensional network structure of matrix is not changed significantly and thus the

tensile fracture strength remains unchanged. However, as the FGE content


99

increases, the crosslink density of epoxy network becomes lower and lower

leading to a large decrease of tensile fracture strength. We can also observe that

the Young’s modulus (at room temperature) of the matrix increases with the FGE

content. All the tested networks are in a glassy state at room temperature and this

variation is similar to the one which was observed by Won [Won1989] for epoxy

networks in which DGEBA was copolymerized with IPD and TMCA

(3,5,5-Trimethylcyclohexylamine), i.e. the corresponding chain extender

(monoamine, f=2). They attributed the variation of Young’s modulus to the

β-relaxation of the TMCA containing networks close to room temperature. In fact, β

amplitude decreased with TMCA content leading to an increase of Young’s

modulus. Since the fracture strength decreases and elastic module increases as a

function of FGE content at the same time, we can deduce that adding FGE leads

to a more brittle matrix. The influence of FGE content on the interfacial mechanical

properties and the self-healing ability will be discussed in later Parts.

(a)


100

(b)

Figure II-16. Evolution of tensile fracture strength (σf) (a) and Young’s modulus (Em) (b) at room

temperature of DGEBA-FGE / IPD matrix as a function of the FGE content (% wt) ( =0.005 min-1).

D. Formation and evaluation of thermally reversible self-healable interphase

In the last two sections, the methods for grafting maleimide groups on the carbon

fiber surface and introducing furan groups in epoxy network were described in

detail. The efficiency of these methods was also confirmed by several analytical

techniques. In this part, the modified carbon surface and epoxy matrix will be

combined to build a thermally reversible self-healable interphase. First of all, the

reversible Diels-Alder reaction used in our system is studied and an appropriate

healing temperature is determined. Then the self-healable interphase is formed

and the micro-droplet debonding test used for measuring the interfacial shear

strength is described.

D.1 Investigation of Diels-Alder reaction

For investigating the process of D-A and retro-D-A reactions, the reaction between

pure BMI and FGE with a stoichiometric ratio of 1:1 in DMSO-d6 solution was


101

studied by nuclear magnetic resonance (1H NMR). The related experimental

parameters about NMR spectroscopy are shown in Annex II. Figure II-17 displays

the two 1H NMR spectra before and after D-A reaction at 60°C for 23 minutes.

(a)

(b) Figure II-17. 1H NMR spectra of D-A reaction in DMSO-d6 solution: (a) at the beginning of reaction

(room temperature) and (b) after reaction at 60 °C for 23 minutes (FGE : BMI = 2:1, mole ratio).


102

It’s obvious that, after reaction at 60°C for 23 minutes, two additional peaks appear

close to 5.3 and 6.6 ppm attributed to the H protons of new formed D-A adducts. In

addition, we can also find that these resonances appear in the form of a

double-peak corresponding to two isomers formed during D-A reaction, i.e.

exo-D-A adduct and endo-D-A adducts (Fig. II-17 (b)).

Figure II-18 1H NMR spectra of D-A reaction (FGE and BMI; molar ratio: 2:1) at different temperatures and reaction times.


103

The D-A reaction kinetic was studied by heating the reaction solution step by step

without interruption. The heating process contains three steps: at room

temperature for 30 minutes, at 60 °C for 3 hours, at 90 °C for 14 hours, and at

110 °C for 30 minutes. During each step, several 1H NMR spectra were recorded

at different times of reaction, shown in Figure II-18. According to this figure, clear

evidences of the reversibility of D-A reactions can be observed with the increase of

temperature, i.e. the thermally reversible character of D-A reaction is proved. From

NMR quantification of the relative resonance peak intensity, the conversion rate of

D-A adduct can be calculated.

Figure II-19. Variation of D-A adduct conversion rate during different temperature steps: X%:

conversion rate of D-A adduct (endo and exo), X1’%: conversion rate of exo structure, X1’’%:

conversion rate of endo structure.

Figure II-19 reports the conversion rate of D-A adducts (endo- and exo-) as a

function of heating time in each temperature step. According to these results, we

can observe that: 1) the D-A reaction can occur at room temperature with a high

reaction rate, i.e. conversion ratio increases rapidly with reaction time; 2) when

temperature was increased to 60°C, the reaction rate starts decreasing, i.e. the


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conversion ratio increases with a relative slow rate, which can be attributed to the

retro-D-A reaction. Because of the retro-D-A reaction, the conversion ratio of D-A

adduct (endo- and exo-) can only reach 67% at 60°C for 3 hours; 3) when the

temperature reaches 90°C, the conversion ratio decreases with a very low rate, i.e.

after 14 hours, about 50% of conversion ratio can be still kept; 4) heating up to

110°C, the D-A adduct conversion ratio decreases rapidly, i.e. keeping at 110°C

for 30 minutes, the D-A conversion ratio decreases to about 20%; 5) comparing

with exo-structure, the endo-structure appears like more unstable, i.e. retro-D-A

reaction is easier to occur with endo-structure, especially at 90°C.

Based on these initial analyses, we can conclude that the D-A reaction

concerned in our interfacial system is thermally reversible. This reaction can occur

rapidly at room temperature. As temperature is between 60 and 90°C, the reaction

is in an equilibrium state and, with the increase of time, this equilibrium state move

to occur retro-D-A reaction. When temperature is higher than 90°C, the retro-D-A

reaction will be accelerated.

Figure II-20. DSC (a) and TGA (b) analyses performed on a stoichiometric mixture BMI : FGE (5

K/min).


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Above conclusions were also confirmed by studying the D-A bond formation by

using Differential Scanning Calorimeter (DSC) and Thermogravimetry analysis

(TGA) for a heating rate of 5 K/min. For mixing BMI and FGE with a stoichiometric

ratio of 1:1, BMI and FGE were dissolved in acetone. After one week at room

temperature, the acetone was removed and a very viscous yellow liquid was

obtained. The DSC and TGA curves from this product are shown in Figure II-20.

On DSC recording during heating, an exothermic platform at about 60°C and

an endothermic peak at about 103°C can be observed. As the temperature

reaches 60°C, the retro-D-A reaction, as an exothermic process, starts happening.

However, when temperature exceeds 90°C, the retro-D-A reaction will be very fast

and the most of the D-A adducts will be separated into BMI and FGE molecules.

Therefore, the endothermic peak can be attributed to evaporation of FGE (boiling

point≈103°C). This analysis is corroborated by the mass loss determined by TGA.

On the TGA trace, the weight loss that takes place between 103 and 200°C is

about 50% of the total weight, which is just similar as the mass ratio of FGE (46%)

in D-A adducts. The second weight-loss plateau can be attributed to the

evaporation or the degradation of BMI molecules.

D.2 Formation of interphase containing thermo-reversible bonds

Figure II-21. Schematic building of a thermo-reversible interphase: D-A thermal reversible bonds,

● irreversible covalent bonds and ◆ Van der Waals force


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Figure II-21 shows the process of the formation of an interphase containing

thermo-reversible bonds generated from the reaction between the maleimide

groups grafted on the carbon fiber surface and the furan groups from FGE which

are with DGEBA to react with isophorone diamine (IPD) for building the

epoxy-amine matrix. Obviously, in this interphase, D-A bonds play not only a role

as healing agent but also as connecting points between BMI-functionalized

carbon fiber and DGEBA-FGE/IPD matrix supplying a necessary interfacial

bonding strength. Thus, this ‘dual role’ of D-A bonds requires that the content of

furan groups in matrix should be appropriate. The best content of furan groups in

the epoxy-amine network should be enough:

i) to react with the maleimide groups on fiber surface to form a strong

interphase;

ii) to be left after debonding for further reaction with maleimide groups during

healing process.

Since the Diels-Alder reaction is a thermo-reversible reaction which tends to form

Diels-Alder adducts below 60°C and return to initial reactants (furan and maleimide)

above 90°C, after curing at high temperature (80°C during 2h +160°C during 2h),

the most part of D-A adducts are in an ‘open’ state. If the temperature decreases to

room temperature rapidly, these separated furan and maleimide groups will not

reconnect in time (Diels-Alder reaction) which will dramatically influence the

interfacial shear strength. Thus, after curing, the samples will be left at 60°C for 2

hours to form more D-A adducts.

As debonding will occur from the micro-droplet debonding test, an interfacial

crack is generated and the droplet slides to a new location on fiber surface. During

the healing treatment, the unreacted furan groups and the ones generated from


107

retro-D-A reaction could react with the maleimide groups to reform D-A reversible

bonds in the fiber-matrix interfacial area. As a consequence, a new interphase

between carbon fiber and epoxy-amine network is generated according to the

formation of new D-A adducts. In Part I, we have also mentioned that D-A covalent

bond is weaker than other bonds in the polymer backbone, so that, when the

material is stressed to its fracture strength, these weakest bonds will be broken

preferentially compared to the other covalent bonds in the network [Thomas2007].

In other words, when D-A based interphase is broken, the D-A adducts will be

completely separated into furan and maleimide groups. Since these groups can

react rapidly at room temperature, this D-A based interphase could also heal itself

at room temperature. Nevertheless, the load-triggered separation of D-A adducts

is more inclined to an unsymmetrical mode (Fig. II-22). For healing this type of

polymer system, an exchange reaction should be carried out. However, the D-A

reaction is in an equilibrium state between 60 and 90°C and the exchange reaction

must be very active. For this reason, we chose a heating at 90 °C for one hour and

a cooling down slowly in the furnace to the room temperature as conditions of

exchange reaction which allow to build again the interphase. In order to ensure a

complete reaction, the interphase-healed composite will be further left at room

temperature for 24 hours before the next interfacial debonding test.

Furthermore, we should also notice that the interfacial fracture mode

mentioned here is very specific, i.e. matrix droplet moves to a new location after

micro-debonding test on fiber surface supplying enough maleimide groups for

rebuilding the fiber/matrix interphase. However, for a real composite material, large

displacements between fiber and matrix never occur in the first steps of fracture,


108

which are concerned by a self-healing solution. In this case, the healing principle is

more likely based on the D-A exchange reaction (Fig. II-22).

Figure II-22. Schematics of the exchange reaction between two D-A adducts in the fiber-matrix

interphase, in the solid state.

D.3 Interphase micromechanics: micro-droplet debonding test

D.3.1 Specimen preparation

The two ends of T700-BMI carbon fiber monofilament were stuck on a metallic

frame using a commercial epoxy glue (®Araldite) after applying a pre-stress (about

400 mN). After mixing the epoxy monomers (DGEBA and FGE) with amine

co-monomer, this reactive mixture was deposited using a thin copper wire to

process micro-droplets of matrix on single fibers (Fig. II-23(a)). Then, the

epoxy-amine droplets were cured following the thermal cycle mentioned previously

(80°C for 2h + 160°C for 2h + 60°C for 2h). As the viscosity of some reactive matrix

systems is very low which makes difficult the deposition of droplet on single fibers,

the matrix was pre-cured at 50 °C for 1 hour to increase the viscosity before

forming the micro-droplets. After curing, the single fibers were cut in segments

about 5 mm in length, at the middle of which there is only one matrix droplet. One

end of the fiber segment was glued on a triangular shape paper to be clamped


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easily in the grips of the tensile machine (Fig. II-23(b)). All the micro-droplets were

observed by optical microscopy to measure the corresponding shape parameters

such as the droplet diameter (Dd), the embedded length (Le), the fiber diameter

(Df), the free length (Lf) between clamping and matrix embedded fiber, the

apparent contact angle (φ), and the true contact angle (θ) (Fig. II-23(b)). Only the

axisymmetric droplets were kept for testing.

(a)

(b) Figiure II-23. Preparation of specimens for micro-debonding test and corresponding geometric

parameters. D.3.2 Process of test

The micro-droplet debonding test was carried out using a tensile testing machine

equipped with a 10N max. load cell having a sensitivity of 0.001N. The

displacement rate and the data collection rate were 0.1 mm/min and 100 Hz,

respectively. The interfacial shear strength was measured by clamping the

micro-droplet between two razor blades (Fig. II-24). The distance between the two

blades can be adjusted for each of the micro-droplets. During debonding test, the


110

upper jig moves upward and the force is recorded as a function of displacement.

When droplet debonding occurs, the tensile force decreases suddenly. Since the

displacement rate is very slow, the maximum tensile force is considered as the

force required for debonding. This force is denoted as F1max. After debonding, the

droplet slides along the fiber with a frictional force denoted as f1.

Figure II-24. Micro-droplet debonding test and setup.

After the first debonding, the debonded micro-droplet samples were heated at

90°C for 1 hour and cooled down slowly in the oven to room temperature and left at

room temperature for 24 hours. This protocol was performed in order to reform the

D-A adducts in the interphase, i.e. in the fractured interfacial zone and to allow the

self-healing of the interphase. In the second cycle, these healed samples were

tested again using the same micromechanical test to repeat above-mentioned

debonding procedure. The debonding force and the sliding friction during the

second cycle were denoted as F2max and f2, respectively. Such a protocol of


111

debonding via micro-droplet test followed by a healing thermal treatment was

repeated several times. Figure II-25 reports the load-displacement trace during

debonding tests repeated several times on a same specimen (the healing

treatment was applied between each debonding test).

Figure II-25. Load-Displacement curves of successive debonding tests performed on the same

micro-droplet prepared from DGEBA-FGE/IPD matrix on T700-BMI carbon fiber.

After testing, the interfacial shear strength (IFSS) was calculated from Function (1),

and the healing efficiency was evaluated by the ratio between the IFSS determined

on the healed interphase (IFSSn+1) and the IFSS of the initial interphase (IFSS1)

(Function (2)). Since the samples for this test are very fragile, any minor mistake

such as breakage of single fiber or vibration during testing etc. will possibly lead to

an incorrect result (or even can’t get any result). Furthermore, the results of

micro-droplet debonding test are generally discrete, the analysis of them requires

a plenty of data points (more than 50 micro-droplets are considered for each


112

interfacial system). For these two reasons, as many as possible samples for each

interphase system should be prepared.

𝐼𝐹𝐼𝐼 = 𝐹𝑚𝑚𝑚𝑆

= 𝐹𝑚𝑚𝑚𝜋𝐷𝑓𝐿𝑒

(1)

𝜂 = 𝐼𝐹𝑆𝑆𝑛+1𝐼𝐹𝑆𝑆1

(2)

E. Conclusion

In this part, the processing of thermally reversible carbon/epoxy interphase was

described in detail, involving in: 1) carbon fiber surface treatments and their

characterization; 2) the matrix composition and its characterization; and 3) building

of the D-A based interphase and the evaluation of interfacial mechanical properties.

Firstly, maleimide groups were grafted on ex-PAN carbon fiber surface through a

three successive steps treatment: i/ nitric acid oxidization, ii/ TEPA amination, and

iii/ BMI grafting. The surface characterization indicated that our treatment method

could effectively introduce maleimide groups on the carbon fiber surface and that

the BMI grafting efficiency deeply depended on the degree of oxidization which

was related to the oxidization time. For this later treatment, the oxidization within

90 minutes could slightly change the surface topography of fiber without

influencing the mechanical properties of carbon fiber. Then, furan groups were

introduced in matrix network by mixing a furan-epoxy monomer (FGE) into epoxy

resin. The characterization results showed that furan groups could be introduced in

the epoxy network from copolymerization with isophorone diamine (and DGEBA).

Furthermore, the ratio of FGE-to-DGEBA could impact the mechanical properties

of matrix and thus influence the interfacial mechanical performance of the following

carbon fiber/ epoxy matrix interphase. At last, the previously modified fiber surface


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and matrix were combined together to interact from a Diels-Alder reaction in order

to design a thermally reversible self-healable interphase. The reaction kinetics of

D-A and retro-D-A reactions were firstly studied and the best healing conditions

were defined, i.e. heating at 90°C for 1 hour and cooling down slowly in the oven

until to the room temperature. Then, processing of micro-droplet debonding test

was described in detail, including sample preparation, testing process, and data

analyses. In the next Parts, the interfacial mechanical properties and the interfacial

self-healing ability will be discussed.


114


115

Part III

Interfacial mechanical properties and self-healing ability of conventional carbon/epoxy interphase A. Introduction ............................................................................................. 117

B. Sample geometrical parameters for micro-droplet debonding test .. 118 B.1 Inter-relationship and significance of embedded length (Le), droplet diameter (Dd), and contact angles (θ and ϕ) ...................................................................... 118 B.2 Significance of free length (Lf) ...................................................................... 122 B.3 Statistics of fiber diameter (Df) ...................................................................... 124

C. Interfacial mechanical properties ......................................................... 124 C.1 Interfacial fracture mode analysis ................................................................. 124 C.2 Analysis of debonding process(es)............................................................... 126

C.2.1 Analysis of the first debonding force ............................................................. 126

C.2.2 Definition of the embedded length (Le) range ................................................ 129

C.2.3 Analysis of the second debonding force ........................................................ 130

C.2.4 Analysis of the frictional force ....................................................................... 131 C.3 Determination of the interfacial shear strength (IFSS).................................. 132

C.3.1 Interfacial shear strength measured during the first debonding (IFSS1) ......... 132

C.3.2 Interfacial shear strength measured during the second debonding (IFSS2) ... 134

C.3.3 Evaluation of the interfacial shear strength of the four reference interphases 135

D. Evaluation of interfacial self-healing ability ......................................... 136

E. Conclusion .............................................................................................. 139


116


117

Part III Interfacial mechanical properties and self-healing ability of conventional carbon/epoxy interphase A. Introduction

In Part II, we have introduced the realization process of D-A reaction based

self-healable carbon/epoxy interphase. The interfacial mechanical properties and

interfacial self-healing ability were proposed to be characterized from the

micro-droplet debonding test. To evaluate the impact of the D-A modification, the

mechanical characterization of conventional carbon/epoxy interphases, composed

from neat and oxidized carbon fibers and non-modified epoxy resin, are taken as

references. This part will report the characterization of the interfacial properties of

these conventional carbon/epoxy interfaces/phases (Table III-1) from the

micro-droplet debonding. The healing ability of these non-modified

interfaces/phases will be also evaluated with the same healing conditions

mentioned in Part II, i.e. heating at 90°C for 1 hour and cooling down slowly to the

room temperature, and left at room temperature for 24 hours.

System Matrix Fiber

I DGEBA (LY556, 𝑡�= 0.15, Huntsman)+IPD

epoxy : amino-hydrogen (1:1)

T700-U II T700-90 III T700-50c IV T700-HNO3 (60min)

Table III-1. Carbon/epoxy interfaces/phases systems used as references. T700-U is the pristine

T700 carbon fiber without any treatment, T700-90 is T700-U after oxidization performed at Toray

Co., T700-50C is an epoxy resin-sized T700-90, and T700-HNO3(60min) is T700-U oxidized in

HNO3 at 115°C for 60 minutes).


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B. Sample geometrical parameters for micro-droplet debonding test

After sample preparation, some shape parameters were measured by optical

microscopy as shown in Part II (Fig. II-23). These parameters, in fact, are not

independent but closely related to each other. The relationship between them will

be discussed and their roles in micro-droplet debonding test will also be defined.

B.1 Inter-relationship and significance of embedded length (Le), droplet diameter (Dd), and contact angles (θ and ϕ)

The contact angle is generally used for evaluating the wettability between a liquid

and a solid surface. The smaller the contact angle is, the better the wettability is.

For a composite system, the wettability of matrix on fiber surface in the first

condition for mechanical performances of the interface/phase. Thanks to optical

microscopy analyses, the apparent (ϕ) and true (θ) contact angles of the cured

epoxy matrix droplet onto the fiber surface were measured to characterize the

wettability of carbon fiber surface (Part II, Fig. II-23). However, in the works of

Yuan [Yuan2013], the contact angles were proposed to be determined from the

matrix droplets before curing. Do the values measured after curing can be used to

evaluate the wettability of epoxy matrix on carbon fiber surface? Figure III-1

displays the average contact angles measured on the four systems taken as

references. Surprisingly, the apparent (ϕ) and true (θ) contact angles show a

different evolution with the fiber surface. According to the information supplied by

the fiber supplier (Toray), the wettability of epoxy matrix on T700-99 (T700-U),

T700-90, and T700-50c increases in turn, i.e. the contact angle of epoxy matrix on

these fibers decreases in turn. Moreover, we have confirmed that, after nitric acid

oxidization, the concentration of polar groups on the fiber surface are increased,


119

which leads to improve the wettability of epoxy matrix on fiber surface. As a

consequence, we propose to use the true contact angle (θ) to characterize the

wettability of matrix on the different fiber surfaces.

The shrinkage of epoxy matrix during curing may change the contact angle

and thus lead to unreliable results. But for the same matrix fully cured in the same

conditions, the shrinkage rate will not change too much. For this reason, the true

contact angle measured after curing is only proposed to qualitatively compare the

wettability of the same matrix on different fiber surfaces.

Figure III-1. Average contact angles between cured matrix droplet and fiber surface.

For the apparent contact angle, its value is also very important for the

micro-droplet debonding test. Chou et al. [Chou1994] have indicated that, for a

given droplet, increasing the distance between the two blades of droplet

debonding test could increase the force component in radial direction and reduce

the force component in axial direction because of the ellipsoidal profile of

micro-droplet. However, the axial force is considered as the main component of


120

shear force applied to the interphase leading to interfacial debonding. In other

words, the larger the distance between the two blades is, the bigger the force

applied for debonding interphase is needed, which leads to an overestimation of

the interfacial shear strength. Similar as the effect of the distance between the

blades, the variation of apparent contact angle can also change the profile of

droplet and thus change the distribution of applied force.

Figure III-2. Force analysis taking place with the micro-droplet debonding test as the apparent

contact angle (ϕ) increases.

Figure III-2 displays the force analysis of micro-droplet debonding with different

apparent contact angles. For a defined interphase (displaying the same interfacial

bonding strength), when the distance between the two blades and the embedded

length is constant, the increase of apparent contact angle will change the profile of

micro-droplet and thus will increase the axial force. In other words, the bigger the

apparent contact angle is, the smaller the applied force is. This conclusion is

similar to one deduced from the works of Choi [Choi2009] and Morlin [Morlin2012]

that artificially controlled the apparent contact angle to about 90° and confirmed

that, with this sample shape, the interfacial debonding is easier to occur.


121

(a) (b)

Figure III-3. Variation of contact angles as a function of the embedded length: (a) the apparent contact angle (ϕ) increases linearly as a function of the embedded length; (2) the true contact angle (θ) is not correlated with the embedded length.

Although the apparent contact angle influences the results of interfacial debonding

test, its value is not independent but related with other parameters. Figure III-3

displays the variation of contact angles as a function of the embedded length.

Obviously, the apparent contact angle increases as a function of the embedded

length. However, for the true contact angle, this tendency does not exist because

that the true contact angle reflects the physiochemical properties of fiber surface

which will not be changed as the increasing of matrix quantity (increase of Le). In

fact, both the embedded length and the apparent contact angle can be considered

as dimensional parameters determined by the volume of matrix droplet.


122

In the process of data analysis, we have also found that the ratio between droplet

diameter (Dd) and embedded length (Le) for the four types of micro-composites is

almost a constant value (Fig. III-4). It means that the interphase of our carbon

fiber/epoxy matrix series does not influence the shape of droplet or, at least, the

impact is not obvious.

Figure III-4. Relationship between droplet diameter (Dd) and embedded length (Le) for the four

different interphases.

According to the analysis above, we can conclude that the shape parameters of Le,

Dd, and ϕ are linearly interlinked and not influenced by the interfacial structure but

determined by the volume of matrix droplet. The effect of these parameters on

interfacial mechanical characterization will be discussed later.

B.2 Significance of free length (Lf)

In Part II, we have underlined that the longer the fiber is, the smaller the fracture

tensile strength is. In fact, the tensile strength of fibers is related to the number of

defects statistically distributed along the fiber, i.e. the tensile strength decreases


123

with the increasing of fiber length. As a consequence, in the process of debonding

test, the tensile strength of the fiber (corresponding to the free length (Lf)) is

required to be higher than the shear strength for interphase debonding; otherwise

the fiber breakage occurs. For ensuring the success of experiment, a large number

of micro-droplet samples with different free lengths (Lf) were tested (Part II, Fig.

II-23). Figure III-5 plots all the free lengths of samples (T700-U/DGEBA) in one

chart. Obviously, the longer the free length is, the easier the fiber breakage occurs.

As a consequence, we propose that the free length should be shorter than 250 µm

to guarantee a high experimental success rate. Since the tensile strength of single

carbon fiber is distributed with an obvious scatter [Pardini2002], 250 µm is only an

approximate upper limit which can roughly guide the tests of other interphases. In

addition, the minimum free length should not be shorter than the thickness of razor

blades, allowing to clamp the droplet.

Figure III-5. Determination of free length to guarantee a high experimental success rate.


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B.3 Statistics of fiber diameter (Df)

During sample preparation, the fiber diameter (Df) was measured by optical

microscopy. The statistical results of fiber diameter are shown in Figure III-6. The

fiber diameter is not a constant value (7µm) but is distributed between 5µm and

10µm. Although this distribution is insignificant for a macroscopic composite, it will

significantly influence the calculation results of embedded area of droplet and thus

the interfacial shear strength. For this reason, the actually measured fiber diameter

will be used for all the further calculations.

Figure III-6. Statistics of fiber diameter.

C. Interfacial mechanical properties

C.1 Interfacial fracture mode analysis

During the micro-droplet dedonding test, four main fracture modes can take place,

including fiber breakage, complete interfacial debonding, cohesive interfacial

debonding, and matrix fracture (Fig. III-7).


125

Figure III-7. SEM images of fracture modes observed after micro-droplet debonding test: (1)

sample before test, (2) complete interfacial debonding, (3) cohesive interfacial debonding, (4)

matrix fracture in bulk.

The appearance of these four rupture modes can be attributed to different stresses

at break, i.e. corresponding forces for a given geometry of the micro-droplet: single

fiber break (Ff), interphase debonding (Fi), and matrix fracture (Fm). The conditions

of fracture are the followings:

Case 1 Ff < Fi & Fm → fiber break,

Case 2 Fi < Ff & Fm → complete interfacial debonding,

Case 3 Fi ≈ Fm < Ff → cohesive interfacial debonding,

Case 4 Fm < Fi & Ff → matrix fracture in bulk.


126

In order to study the interfacial mechanical properties, case 1 and case 4 can be

considered as failed tests and thus the results obtained in these cases can’t be

used for further analysis.

C.2 Analysis of debonding process(es)

C.2.1 Analysis of the first debonding force

Figure III-8 plots the first debonding force (F1max) and the corresponding frictional

force (f1) of the system T700-U/DGEBA as a function of embedded length (Le). The

F1max increases linearly as a function of Le with an obvious scattering. Since the

other three systems have the same trend, only the system T700-U/DGEBA will be

discussed.

Figure III-8. First debonding force and the corresponding frictional force as a function of the

embedded length for T700-U/DGEBA interphase .

The interfacial strength of a composite is provided from not only chemical effects

(covalent bonds connecting fiber surface and matrix network), but also from


127

physical effects (Van der Waals’ interactions and mechanical anchoring of

polymer chains into the pores of fiber surface). In addition, during the interfacial

debonding, a sliding friction triggered by relative displacement between fiber and

matrix takes place. The contributions of these forces will determine the overall

dependence of the first debonding force.

Generally, the interfacial bonding strength is determined by the number and

the type (bond energy) of chemical bonds in interphase, i.e. the more bonds with

high bond energy are formed, the stronger the interphase is. Since the fibers of

each reference interphase are treated in the same conditions and that the matrix is

the same, we can suppose that the number and the type of covalent bonds formed

per unit area of interface is constant. For this reason, the force contributed by

chemical bonds increases linearly with Le (or embedding area, Se). The trend that

the frictional force (f) also increases linearly as a function of the embedded length

(Le) can be observed clearly from Figure III-8. In Part II, we mentioned that

micro-cavities were formed randomly during fiber processing, making possible of

polymer chain-ends or loops to anchor in them. Although the interfacial bonding

strength supplied in this case is random, its contribution to the debonding force is

extremely limited. Therefore, we supposed that it can only provide some

disturbance but not a real influence on the overall tendency of debonding force. By

comparing with bonding strength contributed by the covalent bonds mentioned

previously, the contribution of the Van der Waals’ force appears to be negligible.

The analysis above indicates that F1max linearly increases with Le because that its

main sources (covalent bonds and frictional force) increase linearly as a function of

Le. This trend can be simply displayed as following:


128

F1max = Fp + Fc

= (f+δa+δv)+k1Se

= (kfLe+b) + (δa + δv) + k1πDfLe

= (kf+k1πDf)Le + (b +δa + δv),

where Fp and Fc are the bonding forces due to physical effect and chemical bonds,

respectively; f is the sliding frictional force with f=kfLe+b (kf, b are constants which

can be obtained from curve fitting), δa and δv are the force produced by anchoring

effect and Van der Waals’ force, respectively; k1 is the bonding force resulting from

covalent bonds on per unit area of interface (supposed as a constant), and Se is

the embedded area.

The data scattering can be explained according to four origins: i) the

inhomogeneity of covalent bonds distribution in the interphase. We have assumed

that the fibers are treated uniformly and the matrix is homogeneous. However, it’s

impossible to achieve an absolute uniform fiber surface by oxidization treatment

[Miller1991], leading to an inhomogeneous distribution of covalent bonds at the

interface. For a same embedded length, the number of covalent bonds along the

interface will be different. Furthermore, the distribution of fiber diameter may lead

to the change of interfacial area inducing variation of debonding forces; ii) During

the debonding test, the anchoring forces don’t vary linearly as a function of the

embedded length (Le) but exist randomly; iii) As mentioned by Wu et al. [Wu2014],

the chemical structure of interphase varies gradually from fiber surface to the bulk

matrix which leads to, without doubt, a gradient of mechanical properties. The

fracture takes place in this inhomogeneous interphase leading to the scattering of

the debonding forces; iv) The distance between two clamping blades is another

explanation of the scattering. Heilhecker et al. [Heilhecker2000] mentioned that the


129

vise-angle (Part I, Fig. I-16) that can be controlled by the blades distance

significantly influences the results of micro-droplet debonding test. Although we

have used two micrometers to control the distance between the two blades, it’s

impossible to ensure a constant distance for all the samples because the gap

distance must be adjusted according to the shape of samples; v) The experimental

environment can also play a role on the results accuracy, for example the

fluctuation of environment temperature, air humidity, and vibration, etc.

C.2.2 Definition of the embedded length (Le) range

In the last section, we found that F1max increased linearly as a function of

embedded length (Le). However, the value of Le can’t vary without limitation but in

a specific range. Any sample beyond this range will not be debonded successfully.

This section will determine the range of Le.

Upper limit of the embedded length (Le)

We know that, for achieving a successful interfacial debonding, the interfacial

debonding force (F1max) must be smaller than the tensile fracture force the single

fiber (Ff) between micro-droplet and triangular paper. By taking into account this

condition, the critical embedded length (Lecri) can be predicted with the following

relationship:

F1max ≤ Ff= σf * π(Df/2)2,

where F1max=0.00147Le-0.0107 was obtained from curve fitting (Fig. III-8), σf

(4900MPa) and Df (7μm) are fiber tensile strength and fiber diameter, supplied by

Toray data sheet. The calculation result indicates that Le should be less than


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135μm to ensure a successful interfacial debonding. However, the experimental

results show that some samples with an embedded length longer than 135μm can

still be debonded successfully. The reason for this phenomenon is mainly the

inaccurate parameters (σf and Df) of carbon fiber supplied by Toray. In fact, σf

decreases with the increase of single fiber length and Df distriubuted between 5μm

and 10μm. As a consequence, the experimentally measured values are

proposed to be used for calculatation.

Lower limit of the embedded length (Le)

In theory, the value of Le has no lower limit. However, Figure III-4 displays that the

droplet diameter (Dd) increases linearly as a function of Le, i.e. Dd=0.96Le-18.36.

For a normal micro-droplet, it’s impossible that the Dd is smaller than fiber diameter

(Df). This restriction condition means that Le can’t be reduced without limit.

Calculating with this relationship, the lower limit of Le is confirmed as being about

25μm. But, with this value, it’s very hard to manipulate and requires delicate

equipment. During our experiments, the samples with Le longer than 50μm were

chosen to ease the manipulation.

C.2.3 Analysis of the second debonding force

After the first debonding, all the samples were heated in the conditions that we

have mentioned in Part II, i.e. heating at 90°C for 1 hour and cooled down slowly in

furnace until to the room temperature and left at room temperature for 24 hours.

Then, the samples were tested again and the debonding forces were plotted as a

function of the embedded length. As the four systems display the same variation

tendency, we only reported T700-U/DGEBA for example (Fig. III-9).


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Figure III-9. Plot of the second debonding force and corresponding frictional force as a function of

embedded length.

In theory, after the first debonding, all the covalent bonds existing between fiber

and matrix would be broken and the chain ends anchored in the micro-cavities on

fiber surface would be extracted or broken. The heating step could not make them

reconnect. When the debonding test is repeated after heating, only friction force

and Van der Waals’ interactions can be recovered. For this reason, the second

debonding force decreases significantly in comparison with the first debonding

force and still keeps a linear relation with the embedded length. The composition of

second debonding force can be simply demonstrated as following:

F2max= f +δv

= kfLe+ (b + δv). C.2.4 Analysis of the frictional force

After interphase debonding, the matrix droplet slides along the fiber thus producing

a frictional force (f). Figure III-10 displays the frictional forces of the four reference


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interphases during the first and the second debonding process as a function of the

embedded length.

Figure III-10. Variation of frictional force as a function of embedded length: frictional forces during

the first debonding (1) and the second debonding (2).

Although, from the microscopic point of view, the topography of the fiber surfaces

used for constructing these four interphases are slightly different, these differences

seem not to influence the macroscopic frictional force. The frictional forces for the

considered fiber/matrix systems are very close. Nevertheless, the frictional force

measured during the second debonding is lower than the one determined during

the first debonding. This phenomenon could be attributed to an abrasion effect of

the ‘interface’ during the sliding of the droplet.

C.3 Determination of the interfacial shear strength (IFSS)

C.3.1 Interfacial shear strength measured during the first debonding (IFSS1)

In theory, IFSS is an intrinsic property of interphase. For a given interphase, its

value should be a constant. However, our results show that it is not the case since


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the interfacial shear strength measured during the first debonding (IFSS1)

decreases with the increase of Le (Fig.III-11).

Figure III-11. Dependence of the interfacial shear strength after the first debonding (IFSS1) as a

function of the embedded length (Le) for the four reference interphases.

We have attributed this trend to the variation of load distribution. In the section B.1,

we have studied the relationship between the embedded length (Le) and the

apparent contact angle (ϕ) and we have found that the apparent contact angle

increased with the increasing of the embedded length. In addition, we have also

found that the increasing of the apparent contact angle could change the profile of

micro-droplet and thus increased the component force in axial direction. For these

reasons, we can consider that the increasing of the embedded length can indirectly

increase the ratio of applied force in axial direction which is considered as the main

source of shear force in interphase. In other words, as the embedded length


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increases, more and more proportion of applied force is used for supplying

interfacial shear. As a consequence, for the same interphase, the apparent

interfacial shear strength decreases with the increasing of embedded length.

C.3.2 Interfacial shear strength measured during the second debonding (IFSS2)

Figure III-12. Dependence of the interfacial shear strength after the second debonding (IFSS2) as a

function of the embedded length (Le).

In Figure III-7, displayed in the section C.1, we can easily observe that, after the

first debonding, the embedded length slightly decreases because of the fracture in

Mode I of meniscus. For this reason, the embedded length used for calculating the

second interfacial shear strength (IFSS2) should be measured again before testing.

However, since the samples become very fragile after debonding, further

movement such as clamping again for a new measurement under optical


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microscope will increase the risk of sample damage. In fact, in comparison with

droplet, the dimension of meniscus is very tiny and the embedded length of the first

measurement didn’t count in it (Part II, Fig. II-23). As a consequence, there is no

need to measure again the embedded length for calculating the interfacial shear

strength for the second run. Figure III-12 displays the interfacial shear strength for

a second run of the four reference interphases. Unlike the first interfacial shear

strength, the second interfacial shear strength does not decrease with the increase

of embedded length. It is mainly because that the second debonding only needs to

overcome small forces such as frictional and Van der Waals’ forces. As a

consequence, the distribution of load also becomes unobvious.

C.3.3 Evaluation of the interfacial shear strength of the four reference interphases

Figure III-13. Average interfacial shear strength of four reference interphases after the first

debonding.

Since the values of IFSS obtained by the micro-droplet debonding test are not

constant but decrease with the increase of the embedded length with an obvious

scattering, average interfacial shear strengths were calculated to compare the

interfacial strength of the four different interphases. Figure III-13 displays the


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average interfacial shear strength of the four reference interphases after the first

debonding.

Obviously, the shear strengths of the interphases constructed with treated

fibers (T700-90, T700-50c, and T700-HNO3 (60min)) are higher than that

constructed with non-treated fiber (T700-U). The fiber surface treatment, on the

one hand, increases the polar groups onto the fiber surface which allows to

generate covalent bonds with the epoxy network matrix; on the other hand, the

surface treatment changes the fiber surface topography, especially after nitric acid

oxidization treatment, allowing more polymer chain to be anchored on the fiber

surface. According all these changes, the fiber surface treatment increases the

interfacial shear strength.

D. Evaluation of interfacial self-healing ability

In the forward sections, we have mentioned that the interfacial strength is mainly

governed by four types of interactions, i) covalent bonds, ii) friction, iii) anchoring,

and iv) Van der Waals’ interactions. When the interphase is debonded during the

first run, the covalent bonds and the anchoring will not exist anymore. However, if

we continue to test the mechanical properties of the debonded ‘interface’, the

results are not zero because of friction and Van der Waals’ interactions.

Figure III-14 displays the interfacial self-healing efficiencies (𝜂) of the four

reference interphases, which were obtained by calculating the ratio of the first and

the second interfacial shear strength (Part II, Function (2)). All the composite

systems possess more or less interfacial self-healing abilities and these abilities

seem to be related with fiber surface structure (chemical and physical). The


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relationship between the fiber surface structure and the interfacial self-healing

efficiency will be explained subsequently.

Figure III-14. Interfacial self-healing efficiency of reference interphases.

Since T700-U did not undergo any treatment, neither cavities nor functional groups

should be present on its surface. Associated with the epoxy matrix, the interfacial

strength will be mainly governed by friction and Van der Waals’ interactions

leading to a relatively weak interphase (Fig. III-13). However, both of these

interactions are recoverable, which contribute to a relatively higher healing

efficiency. After the oxidization treatment, the functional groups on fiber surface

will be increased significantly, with which the number of covalent bonds will be

dominant in interphase. These covalent bonds can obviously improve the

interfacial strength but can’t be recovered. As a consequence, the interphase

constructed with T700-90 is relatively strong but has a relatively low healing

efficiency. For the interphase based on T700-50c, the new formed ‘interface’ after

the first debonding is actually between the epoxy sizing of fiber and epoxy bulk.

The free chain ends of epoxy can interpenetrate and inter-diffuse after heating and


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thus leading to a relatively high healing efficiency. We have confirmed that cavities

can be formed after oxidization in nitric acid, which allow some polymer chains to

enter during heating. When these polymer chains were firmly stuck in the cavities,

the situation is similar to the one for which covalent bonds exist between fiber and

matrix bulk. That is the reason why the healing efficiency of the interphase

composed by T700-HNO3 (60min) is higher than other interphase systems even

though the number of irreversible covalent bonds is dominant during the first

debonding (the same case as the interphase composed by T700-90).

In fact, the above analysis only highlights the key factors for each interphase,

which influence the interfacial healing efficiency. Besides these, friction and Van

der Waals’ interactions could also contribute to the healing efficiency of each

interphase system. However, friction contributes to the interfacial strength only

during sliding of the droplet on fiber. Thus, the interfacial healing ability wouldn’t

consider the contribution of the frictional force. As a consequence, the interfacial

healing efficiency could be calculated as following:

𝜂∗ = 𝐼𝐹𝑆𝑆𝑛+1𝐼𝐹𝑆𝑆1

= 𝐹𝑛+1−𝑓𝑛+1𝐹1−𝑓1

(n=1, 2, 3…), (1)

where 𝜂* is the interfacial healing efficiency without taking into account the friction

contribution. F and f are debonding and corresponding frictional forces,

respectively. Figure III-15 shows the healing efficiency of the reference

interphases excluding the friction effect. Obviously, the conventional carbon fiber /

epoxy matrix interphases don’t possess the ability of self-healing.


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Figure III-15. Interfacial self-healing efficiency of reference interphases (without frictional effect

taken into account).

E. Conclusion

In this part, the basic characteristics of micro-droplet debonding test were

discussed by studying the interfacial mechanical properties of reference

interphases: i) neat ex-PAN carbon fiber without any treatment (T700-U); ii)

commercially oxidized carbon fiber (T700-90); iii) commercially epoxy-sized

carbon fiber (T700-50c); iv) HNO3 oxidized carbon fiber (T700-HNO3) . The

discussion is carried out in three aspects, including the relationship and the

influence of the micro-droplet shape parameters, the interfacial strength of

reference interphases, and the self-healing ability of reference interphases.

In the first sub-part, the significance and influence of the contact angles (true

(θ) and apparent (ϕ)) formed between fiber and cured matrix droplet, on the

micro-droplet debonding test are discussed. The true contact angle (θ) can be

used to qualitatively compare the wettability of epoxy matrix on different carbon

fiber surfaces. However, the apparent contact angle (ϕ) can be considered as a

dimensional parameter and influence the testing results because of the change of


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droplet profile. Furthermore, we have also found that the dimensional

characteristics of the matrix micro-droplet (Le, Dd, and ϕ) vary linearly with each

other. These correlations are not dependent on the structure of interphase

(physical and chemical) but on the volume of micro-droplet. From a fine analysis of

the various geometrical practical considerations, the conditions of micro-droplet

processing and micro-mechanical testing were defined in order to compare the

different interfaces/phases considered in this work.

In the second sub-part, the fracture modes observed after micro-droplet

debonding were identified: fiber break, complete interfacial debonding, cohesive

interfacial debonding (cohesive failure close to the interface), and fracture in the

bulk matrix. Only the complete and the cohesive interfacial debonding can be used

for analyzing the interfacial mechanical properties. Then, we analyzed the

debonding force and indicated that both the first and the second interfacial

debonding forces increase linearly as a function of embedded length with an

obvious scattering. The reasons for data scattering were also discussed in detail.

In addition, we also found that the frictional force caused by relative movement

between fiber surface and matrix increases with the embedded length and

decreases with the debonding times because of ‘interface’ abrasion. After

calculating, we have found that the interfacial shear strength measured by droplet

debonding test is not a constant but decreases with the increasing of embedded

length which can be attributed to the variation of load distribution. At last, the

interfacial shear strengths of the four reference interphases were compared from

the average interfacial shear strength. With this comparison, we have concluded

that the interphases based on treated carbon fibers possess a better interfacial


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bonding strength, which can be attributed to the formation of chemical and physical

bonds between fiber surface and matrix.

In the last sub-part, the healing abilities of the reference interphases were

studied. The results demonstrated that the conventional carbon/epoxy interphase

has the ability to recover more or less the initial interfacial shear strength and the

recovery ability is related to the interfacial structures. However, this recovery ability

can’t be considered as an interfacial self-healing because that it’s mainly due to

the passive frictional force.

In the next part, the mechanical properties and healing abilities of D-A

modified interphase will be studied. The feasibility of interfacial self-healing based

on D-A reaction will be confirmed and the best fiber surface treatment process and

matrix component will be also discussed.


142


143

Part IV

Interfacial mechanical properties and self-healing ability of D-A modified interphase A. Influence of carbon fiber surface treatment on interfacial mechanical performances and interfacial self-healing ability of the D-A modified interphase .................................................................................................... 145

A.1 Introduction ................................................................................................... 145 A.2 Wettability of furan-modified epoxy matrix on different fiber surfaces .......... 147 A.3 Analysis of interfacial shear strength ............................................................ 148 A.4 Evaluation of interfacial self-healing ability ................................................... 150

A.4.1 Feasibility of interfacial self-healing ............................................................... 150

A.4.2 Interfacial self-healing efficiency.................................................................... 155

A.4.3 Comparison healing abilities of the different D-A modified interphases ......... 157

A.4.4 Repeatability of interfacial self-healing .......................................................... 160 A.5 Conclusion .................................................................................................... 162

B. Influence of FGE content on D-A modified interphase properties .... 164 B.1 Introduction ................................................................................................... 164 B.2 Influence of FGE content on the interfacial mechanical properties ............... 165

B.2.1 Influence on the interfacial bonding strength ................................................. 165

B.2.2 Influence of FGE content on the interfacial toughness .................................. 166 B.3 Influence of the FGE content on the interfacial self-healing efficiency.......... 168

B.3.1 Characterization of the interfacial self-healing efficiency based on IFSS ....... 168

B.3.2 Characterization of interfacial self-healing efficiency based on Gic ................ 170 B.4 Conclusion .................................................................................................... 173


144


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Part IV Interfacial mechanical properties and self-healing ability of D-A modified interphase

This part will analyze the results of micromechanical tests in order to define the

best conditions for designing a thermally self-healable carbon/epoxy interphase,

which, on the one hand, possesses the ability of self-healing and, on the other

hand, displays a good overall mechanical performance. The discussion will be

carried on considering two aspects, including: (1) the influence of carbon fiber

surface treatment on the interfacial mechanical performance and the interfacial

self-healing ability; (2) the influence of the matrix composition on the interfacial

mechanical performance and the interfacial self-healing ability.

A. Influence of carbon fiber surface treatment on interfacial mechanical performances and interfacial self-healing ability of the D-A modified interphase

A.1 Introduction

We have shown that D-A reversible covalent bonds could be successfully

introduced into carbon/epoxy interphase by fiber surface grafting and epoxy matrix

modification. This section will evaluate the mechanical properties and self-healing

ability of this type of interphase by comparing the results of successive debonding

tests performed on micro-droplets. In the last part, we have underlined that the

number of chemical bonds between fiber surface and matrix significantly

influenced the mechanical properties of interphase. For the D-A covalent bonds

based interphase, this conclusion is still available. Furthermore, if the interfacial


146

self-healing based on D-A reaction is feasible, the number of new formed D-A

bonds will directly determine the interfacial healing efficiency. Since carbon fiber

surface is inert, the grafting efficiency of BMI is very limited and thus the number of

maleimide groups on fiber surface will be a key factor that will determine the

number of D-A bonds in interphase. However, the number of maleimide groups on

fiber surface is closely related to the oxidization time before grafting with BMI. In

other words, the fiber oxidization time will determine the number of D-A covalent

bonds and also the interfacial performances (mechanical properties and

self-healing ability). To check it, a series of interfacial systems will be designed by

combining epoxy matrix with a serial of BMI-grafted carbon fibers which have been

oxidized in nitric acid for different times before grafting. The different interfacial

systems are reported in Table IV-1. The influence of oxidization time on the

interfacial shear strength and the self-healing ability will be mainly discussed and

the appropriate oxidization time will also be proposed. Furthermore, each

interfacial system will undergo more than four debonding-healing cycles, for which

the repeatability of self-healing will be confirmed. Here, we choose a relatively high

DGEBA/FGE (6:4 by wt.%) to ensure that all the maleimide groups on fiber surface

will be consumed to form D-A bonds. The best weight ratio of FGE will be

discussed in the next section.

System Matrix Fiber

I DGEBA (LY556, 𝑡�= 0.15, Huntsman)/FGE+IPD

DGEBA/FGE weight fraction: 6:4 Epoxy : H-amine = 1:1

T700-U II T700-BMI(30min) III T700-BMI(60min) IV T700-BMI(90min)

Table IV-1. D-A modified interphase systems with a serial of carbon fibers oxidized for different

times.


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A.2 Wettability of furan-modified epoxy matrix on different fiber surfaces

In Part III, we have proposed that the true contact angle (θ) measured after curing

can be used to qualitatively compare the wettability of the same matrix system on

different fiber surfaces. Figure IV-1 displays the true contact angles of the four D-A

modified interphases. In theory, there are many polar functional groups on BMI

molecule such carbonyl and carbon-carbon double bonds, the contact angle on the

BMI-modified interface should be decreased, i.e. leading to a good wettability.

However, our results show the opposite. This phenomenon can be attributed to: 1)

the presence of numerous polar groups on BMI molecule. Their polarity can screen

each other and thus lead to a non-polarity; 2) as the epoxy matrix is a viscous

liquid, the increase of fiber surface roughness caused by oxidization and grafting

process can also influence the measurement of contact angle.

Figure IV-1. Determination of the true contact angles measured between epoxy matrix and carbon

fiber oxidized or not for varying times: 30, 60, and 90 minutes.


148

The change of wettability is mainly influenced by the physical absorption of the

polymer matrix on the fiber surface, i.e. Van der Waals’ forces. In the D-A modified

interfacial systems, the contribution of this type of interactions to interfacial shear

strength is very limited. Therefore, it will not influence the overall trend of interfacial

mechanical properties.

A.3 Analysis of interfacial shear strength

In Part II, we have confirmed that, within 90 minutes, the number of acidic

functional groups increased as a function of oxidization time. According to this

conclusion, the number of maleimide groups grafted on fiber surface should exhibit

the same trend as well as the interfacial shear strength. However, the reported

results do not show this tendency (Fig. IV-2).

Figure IV-2. Average interfacial shear strength of the D-A modified interphases during the first

debonding.

Within 60 minutes, the average interfacial shear strength at the first debonding

increases as a function of oxidization time. After 90 minutes, this value decreases.


149

This dependence can be attributed to the surface cavities formed after oxidization.

As mentioned previously, for 90 minutes of oxidization treatment, the deeper

surface layer of fiber is etched and interconnected pores or internal voids and

cracks can be created in the sub-surface. In these void structures, a large number

of acidic groups could be expected. However, these functional groups can’t be

used to graft with maleimide but hidden under the TEPA grafting layers. As a

consequence, the interfacial shear strength doesn’t increase with the increasing of

acidic functional groups.

In fact, the XPS analyses indicated that the more acidic functional groups on

fiber surface, the lower the TEPA grafting efficiency is (Part II, Table II-4). This

phenomenon can be attributed to the ‘back-biting’ of TEPA molecule on the fiber

surface. In other words, the fiber surface oxidized for 30 minutes displays the

higher grafting ratio of amine groups and, as a consequence, higher grafting ratio

of maleimide groups. Nevertheless, the interfacial shear strength of the

corresponding interphase is not the highest. We could explain this surprising result

by the schematic drawing, shown in Figure IV-3.

Figure IV-3. Schematic drawing showing the influence of TEPA molecule ‘back-biting’ on interfacial

bonding.


150

Without back-biting’, there is only one connection formed between fiber surface

and TEPA molecule (Fig. IV-3(a)). Even if all the amino-hydrogens can react with

BMI to graft maleimide groups and further to form D-A covalent bonds with epoxy

network, the interfacial shear strength will be only dependent on one covalent bond.

However, if there is ‘back-biting’, the connection points between the fiber surface

and the bulk matrix will be increased significantly, i.e. lead to a stronger interphase

(Fig. IV-3(b)).

A.4 Evaluation of interfacial self-healing ability

A.4.1 Feasibility of interfacial self-healing

In the last Part, we proposed to represent the interfacial shear strength by a simple

model expressed in terms of force balance. This model applied to the D-A modified

interphase can be expressed as following:

F1max = Fp + Fc

= (f1+δa+δv)+(Fc’+ Fc

”)

= (kf1Le+b) + (δa + δv) + (k’πDfLe+ k”πDfLe)

= (kf1+k’πDf+ k”πDf)Le + (b +δa + δv),

where Fc, Fp, Fc’, and Fc

” are the bonding forces corresponding to the chemical

bonds, physical bonds, non-reversible covalent bonds, and D-A reversible

covalent bonds, respectively; f, δa, and δv are sliding friction force (f=kf1Le+b; kf1,

b are constants which are obtained from curve fitting), the force produced from

anchoring effect and Van der Waals’ forces, respectively; k’ and k” are the bonding

forces per unit area of interphase, owing to non-reversible covalent bonds and D-A

reversible covalent bonds, respectively. Figure IV-4 plots the first debonding force


151

for the interface systems as a function of the embedded length. As expected, the

first debonding force increases linearly as a funtion of Le.


152

Figure IV-4. First debonding force and the corresponding frictional force as a function of the

embedded length for various interfacial systems.

When the debonded interphase undergoes a heating process, the equations for

the second debonding force could be expressed as following:

F2max - f2= k”πDfLe + δv .


153

In last part, we have mentioned that the frictional force can’t be considered as the

healing ability and thus the term of frictional force (f2) is moved to the left of

equation. According to this equation, if the D-A covalent bonds were formed during

heating, the value of (F2max-f2) will display a linear relationship with Le, otherwise

this value will be distributed randomly with a very low value. Figure IV-5 plots the

value of (F2max -f2) as a function of the embedded length.


154

Figure IV-5. Second debonding force (excluding frictional force) as a function of the embedded

length for the different interfacial systems.

Obviously, for the D-A modified interphases, (F2max-f2) increases linearly as a

function of embedded length. However, for the reference interphase (based on

T700-U), this value is roughly constant. This result confirms that the interfacial

self-healing based on thermally reversible Diels-Alder reactions is feasible.


155

A.4.2 Interfacial self-healing efficiency

Figure IV-6 displays the first interfacial self-healing efficiency as a function of the

embedded length. The frictional forces have been excluded for calculation.

Obviously, the interfacial healing efficiencies of the four systems decrease with the

increasing of embedded length, especially the D-A modified interphases.


156

Figure IV-6. Interfacial self-healing efficiency of D-A modified interphases without frictional effect.

Both values of F2max-f2 and F1max-f1 used for calculating the first healing efficiency

increase as a function of embedded length. The only reason why their ratio

decreases as a function of embedded length is that their variation is not the same.

Before the first debonding, the contact between fiber surface and matrix is very

close because that the liquid state epoxy matrix can infiltrate the whole fiber

surface before curing. For this reason, the value of F1max-f1 will keep the same


157

variation with embedded length, i.e. the value of F1max-f1 increases linearly with the

increasing of embedded length. However, before the second debonding, the matrix

has been cured and the additional heating will not significantly improve the contact

between fiber surface and bulk matrix droplet. Furthermore, we have confirmed

that the fiber diameter is not uniform but is distributed in a range even though the

same fiber was measured at different locations. In this case, the micro-droplet with

a longer embedded length is easier to be ‘built on stilts’, and thus reforming

relatively less D-A bonds (Fig. IV-7). In other words, the increasing rate of F2max-f2

will decrease with the increasing of embedded length because of the ‘stilts’ effect.

For the reference interphase (composed by T700-U), the recovery of interfacial

bonding strength mainly depends on Van der Waals’ forces which are very weak

interactions between fiber surface and bulk matrix. Therefore, the change of

F2max-f2 caused by the increasing of embedded length is not notable (Fig. IV-5).

Figure IV-7. Schematic drawing illustrating the difference of fiber/matrix contact modes: before the

first debonding, fiber surface and matrix closely contact with each other; before the second

debonding, the droplet with longer embedded length is easier to be ‘built on stilts’.

A.4.3 Comparison healing abilities of the different D-A modified interphases

Figure IV-8 displays the average interfacial self-healing efficiencies of the D-A

modified and reference interphases. Obviously, the average healing efficiency of


158

D-A modified interphases are higher than that of reference interphase, which

confirms the success of D-A based interfacial self-healing. Among the three D-A

modified interphases, the best result (75%) was achieved by using T700-BMI

which has been oxidized for 60 minutes during the first treatment step.

Figure IV-8. Average interfacial self-healing efficiency of D-A modified interphases (without

frictional effect).

This phenomenon can be explained by the state of micro-droplet before the first

and the second debondings. Before the first debonding, the fiber surface and the

matrix are in close contact and the interfacial shear strength mainly depends on

the fiber surface treatment. However, after the first debonding, the relative

displacement between fiber surface and matrix droplet changes the original

surface structure of fiber due to friction. However, the frictional resistance for the

three types of fiber surfaces is not the same. For example, we have mentioned that,

after grafting TEPA and BMI on fiber surface which has been oxidized in relatively

short time, a ‘tree-like’ structure would be formed (Fig. IV-3). This type of surface

structure is easier to be abraded during relative displacement between fiber and


159

matrix droplet. Thus, as a consequence, the number of available maleimide groups

on fiber surface which can be further used to reform the interphase decreases. For

the fiber oxidized for 90 minutes, its surface has been over oxidized and become

loose because of the cavities in sub-surface. Debonding the interphase formed by

this type of fiber, the propagation of crack is more likely between the grafted layer

and the graphite fiber surface or even peels the fiber surface. The peeled piece will

follow the droplet to a new location on fiber surface and will further hinder the

formation of new D-A bonds.

(a)

(b) Figure IV-9. Schematic drawing of fiber surface peeling (a) and typical SEM image of T700-BMI

(90min) fiber surface after the first debonding (b).


160

Figure IV-9 displays the proposed scenario of fiber surface peeling and the typical

SEM image of T700-BMI(90min) fiber surface after the first debonding, showing

surface peeling. Nevertheless, for the fiber oxidized for 60 minutes, both cases

mentioned previously do not exist or are not significant. As a consequence, the

interphase formed can get a best healing efficiency.

A.4.4 Repeatability of interfacial self-healing

Park et al. [Park2010] have mentioned that the Diels-Alder bonds are weaker than

the other covalent bonds in a polymer backbone. They should be broken

preferentially under stress. When the broken interphases are heated, all the

separated D-A bonds will be reconnected. So, theoretically, these D-A bonds

existing in the interphase can contribute to heal the interfacial fracture several

times without any reduction of healing efficiency. Nevertheless, the results shown

in the Figure IV-10 indicate that the healing efficiency is reduced step-by-step as

the debonding-healing cycle number is increasing.

Figure IV-10. Evolution of the self-healing efficiency as a function of debonding-healing cycles for

DGEBA/FGE (6:4) epoxy matrix/T700-BMI carbon fiber (prepared from T700-HNO3 carbon fiber

oxidized for 60 min).


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This evolution can be attributed to the reduction of furan groups on the matrix side.

During the debonding, friction, uneven forces and displacement rate etc. will lead

to an asymmetrical separation of D-A adducts. In this case, some of the furan

groups will be left, in the form of D-A adduct on fiber surface after extraction from

the matrix (cohesive debonding). Near the newly formed ‘interface’ (before

healing), the concentration of residual furan groups that will further be used for

reforming D-A adducts is reduced (Fig. IV-11). Furthermore, repeated friction can

also increase the gap between fiber surface and matrix droplet which is

unfavorable for further interphase rebuilding.

Figure IV-11. Interphase debonding and self-healing processes: ● irreversible covalent bonds ◆

Van der Waals’ force

Figure IV-12 shows SEM micrographs of the carbon fiber surface of a specimen

after several debonding-healing cycles. Obviously, after each debonding, there are

some pieces of epoxy matrix that remain on the fiber surface and that can contain

furan groups. Until the fourth debonding, such type of cohesive fracture becomes

unobvious.


162

Figure IV-12. SEM images of fiber surface after 4 debonding-healing cycles: I, II, III, IV are the

partial enlarged views of fiber surface after each debonding.

A.5 Conclusion

In this section, the interfacial mechanical properties and self-healing abilities of the

D-A modified interphases were discussed in detail and some conclusions were

obtained as following:

1) Although grafting BMI on fiber surface worsens the wettability of epoxy

matrix, the interfacial shear strength is increased because that D-A covalent

bonds are formed between fiber surface and matrix. Forming the interphase

with T700-BMI which has been oxidized for 60 minutes during the first

treatment step can get the highest interfacial shear strength (89MPa).

2) By comparing the evolution of the second debonding forces, the feasibility

of interfacial self-healing based on Diels-Alder reaction was confirmed. The

interfacial self-healing efficiency of D-A modified interphase is not a


163

constant but decreases with the increasing of embedded length, caused by

the ‘built on stilts’ of droplet.

3) The highest healing efficiency (75%) is achieved by using the carbon fiber

T700-BMI which was previously oxidized in HNO3 for 60 minutes, i.e.

allowing to introduce a higher concentration of D-A adducts into the

interphase. The interphase of this system can be healed several times and

the healing efficiency is reduced step-by-step as the debonding-healing

cycle number is increasing. The reason for this variation can be attributed to

the abrasion of ‘interface’ between fiber surface and matrix droplet.

For this kind of self-healing interphase, the interfacial mechanical properties and

healing abilities are mainly determined by the fiber surface treatment process.

However, for a composite material, the overall properties are also very important,

which requires to focus not only on the fiber and its surface but also on the

properties of matrix. In fact, in our study, the mechanical properties of matrix are

mainly determined by the amount of FGE which acts as chain extender between

crosslinks. 40% of FGE (weight fraction) in matrix is obviously excessive, which

leads to a relatively low glass transition temperature (Tg≈63°C) and mechanical

properties (Part II, Fig. II-15 and Fig. II-16). In the next section, the weight fraction

of FGE will be studied by considering series of interphases composed by

T700-BMI (oxidized for 60 minutes during the first treatment step) and

epoxy-amine matrix systems with different FGE weight fraction.


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B. Influence of FGE content on D-A modified interphase properties

B.1 Introduction

Thermo-reversible Diels-Alder (D-A) bonds formed between maleimide and furan

groups have been used to generate an interphase between carbon fiber surface

and an epoxy matrix leading to the ability of interfacial self-healing in carbon/epoxy

composite materials. However, for ensuring all the maleimide groups on fiber

surface could be consumed to form D-A bonds, an excessive content of FGE was

added in matrix (DGEBA:FGE, weight fraction of 6:4). As FGE acts as chain

extender in epoxy network, such FGE content significantly weaken the thermal and

mechanical properties of matrix and thus the overall properties of the composites.

In this section, the influence of FGE content on the interfacial mechanical

properties and the self-healing abilities will be studied by comparing series of

interphases combining carbon fiber and epoxy matrix containing different weight

fractions of FGE. The compositions of these interphases are shown in Table IV-2.

System

Matrix DGEBA

(D.E.R.TM 331, 𝑡�= 0.15, Dow) /FGE+IPD

Epoxy : H-amine = 1:1

Fiber

DGEBA/FGE weight fraction

T700-BMI(60min)

Oxidized for 60 minutes during the first treatment

step

I 10:0 II 9:1 III 8:2 IV 7:3 V 6:4 VI 5:5

Table IV-2. Interfacial systems used for studying the influence of FGE content on interfacial

performance of D-A based self-healable interphases.


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The study of interfacial mechanical properties will mainly include the interfacial

shear strength and the interphase toughness. The interfacial self-healing ability will

be evaluated from the analysis of the successive debonding forces and the nature

of the interphase fracture via its critical energy release rate (Gic). By considering

both the interfacial mechanical properties and the healing ability, the best fraction

of FGE for modifying the matrix will be proposed. For such a formulation, the

composite system will have not only a good interfacial self-healing ability but also

will maintain relatively high mechanical performances.

Figure IV-13. Evolution of the first debonding force as a function of the embedded length for

different weight fractures of FGE introduced in the epoxy matrix.

B.2 Influence of FGE content on the interfacial mechanical properties

B.2.1 Influence on the interfacial bonding strength

In theory, as the FGE content in matrix increases, the number of D-A adducts

formed in interphase should increase gradually until all the maleimide groups on

fiber surface are reacted. During this process, the interfacial bonding strength


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which is determined by the number of covalent bonds should follow the same

dependence, i.e. increasing until to a constant value.

Figure IV-13 reports the first interfacial debonding force as a function of the

embedded length for various FGE contents in the epoxy matrix. It’s obvious that

the debonding force increases with the FGE content. However, as the FGE content

is higher than 40%, the debonding force slightly increases with FGE content.

Therefore, 40% wt. can be considered as a critical value, for which all the

maleimide groups on fiber surface can react to form D-A bonds. The average

interfacial shear strength (IFSS) calculated from the first debonding step is shown

in Figure IV-14. It clearly shows that the IFSS reaches a plateau value for 40% wt.

of FGE.

Figure IV-14. Dependence of the interfacial shear strength as a function of the embedded length for

different contents (by wt.) of FGE introduced in the epoxy matrix (first debonding step).

B.2.2 Influence of FGE content on the interfacial toughness

As the FGE monomer modified the crosslink density of the bulk matrix network, it

also changes the crosslink density of the interfacial layer, i.e. the epoxy matrix


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close to the fiber surface (Fig. IV-15). In fact, decreasing the crosslink density

leads to increase the fracture toughness of epoxy networks [Won1989]. As a

consequence, the interphase layer in FGE-containing systems will display a higher

toughness.

Figure IV-15. Schematic drawing of the change of crosslink density in the interfacial zone with

increasing of the FGE content. The FGE weight fraction in the epoxy matrix A is lower than that in

matrix B.

For characterizing the interfacial toughness of the considered systems, the

elongation at fracture of interphase, i.e. the ratio between the displacement of

micro-droplet from the beginning of loading to debonding and the embedded

length was measured, which can be qualitatively considered as an ‘interfacial

shear strain’ and could be related to the interfacial toughness. The average values

reported in Figure IV-16 show that the interfacial toughness increases with the

increasing of FGE content.

Here, we should notice that although the interfacial strength and toughness

increase as a function of the FGE content, other physical properties of the bulk

epoxy matrix which determine the overall properties of the composite material

decrease significantly, such as glass transition temperature and fracture strength

(Part II, Fig. II-15 and Fig.II-16). As a consequence, the optimal FGE content for

D-A based self-healable interphase can’t be confirmed so far.


168

Figure IV-16. Average elongation at fracture of interphase for debonding as a function of FGE

weight fraction.

B.3 Influence of the FGE content on the interfacial self-healing efficiency

B.3.1 Characterization of the interfacial self-healing efficiency based on IFSS

According to the calculation method mentioned in Part III, the average interfacial

self-healing efficiencies excluding frictional effect were calculated. The results are

reported in Figure IV-17.

Figure IV-17. Interfacial self-healing efficiency (excluding frictional effect) as a function of the FGE

weight fraction.


169

In this figure, a threshold value close to 20 wt% of FGE can be clearly evidenced.

When the content of FGE is lower than 20% in the epoxy matrix, a relatively high

healing efficiency can be obtained (about 80%). However, as the FGE content

increases up to 20%, the healing efficiency gradually decreases towards a much

lower value. This effect can be explained by the interfacial bonding mode and

fracture mode. When the FGE content is low (less than 20%), the interfacial

strength is mainly related to the physical bonds, such as Van der Waals’ force, and

friction. These contributions of physical bonds can be recovered after debonding,

thereby getting a higher healing efficiency. However, comparing with covalent

bonds, the contributions of physical bonds at the fiber/matrix interphase are

negligible. As a consequence, the interfacial strength is low (Fig. IV-14). With the

increasing of FGE, more and more D-A covalent bonds are formed in the

interphase and could contribute to increase the interfacial strength. These changes

lead to an unstable crack propagation in the interphase during debonding, i.e.

interphase is no longer a weak link and cracks could propagate into the bulk matrix.

After debonding, some small pieces of epoxy matrix remain on the fiber surface.

This phenomenon becomes more and more important with increasing of the FGE

content (Fig. IV-18). In the matrix, the number of furan groups involved in D-A

adducts is large. As a consequence, the furan group content close to the new

formed interface decreases, which should further be used to reform D-A adducts.

Furthermore, due to the large quantity of matrix left on the fiber surface as

fragments, a close contact between the fiber surface and the matrix required for

healing (from reaction or exchange of furan and maleimide groups) will be difficult.

When the interphase is seriously damaged, no D-A bonds could be reformed by

heating and the healing efficiency is low.


170

Figure IV-18. Typical SEM images of fiber surface after the first debonding.

B.3.2 Characterization of interfacial self-healing efficiency based on Gic

Scheer et al. [Scheer1995] demonstrated that the determination of IFSS by

considering a constant shear stress along the embedded length (Part II, Function

(1)) is not correct. As a consequence, they proposed to determine the energy

release rate for debonding, denoted as critical energy release rate (Gic) to

characterize the interfacial properties. The relative shear-lag calculation function is

shown as following, which is reported in detail in Annex II.


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Table IV-3 summarizes the fiber and matrix parameters required in this

micromechanical model, including fiber diameter (Df), fiber tensile modulus (Ef),

fiber axial thermal expansion coefficient (αf), and matrix thermal expansion

coefficient (αm). According to Won [Won1989], the values of thermal expansion

coefficient of the matrix do not change strongly with the crosslink density of the

epoxy matrix. In addition, Scheer [Scheer1995] has also indicated that the precise

values of these properties have not a large influence on Gic. For this reasons, the

same value of αm was considered for all of the epoxy matrix systems.

Parameters Value Reference

Df 7μm Toray Data sheet Ef 2 3 0 G P a Toray Data sheet αf -0.4 ~ -1.0 x10-6/K, Toray Data sheet αm 171 x10-6/K Won1989

Table IV-3. Physical characteristics of carbon fiber and epoxy matrix required for the application of

the shear-lag model.

Figure IV-19 displays the calculated critical energy release rate of the first

debonding (Gic1) and the ratio between the second and the first energy release rate

(Gic2/Gic1) which can be considered as the interfacial self-healing efficiency. The

value of Gic1 and Gic2/Gic1 exhibit almost the same dependence as the interfacial

shear strength and interfacial self-healing efficiency calculated previously.

According to these results, we propose that the optimum FGE content in the

epoxy matrix should be between 20% and 30%, for which the interphase can not

only get the best healing efficiency but also maintain a relatively high interfacial

mechanical performance. Furthermore, with this FGE content, the thermal and


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mechanical properties of matrix also maintain relatively high values, i.e. average

fracture strength (σf) and glass transition temperature (Tg) are 70MPa and 100°C,

respectively (Part II, Fig. II-15 and Fig. II-16).

Figure IV-19. Dependence of the critical energy release rate at the first debonding (Gic1) and the

ratio between the second and the first energy release rate (Gic2/Gic1) as a function of FGE content.

One can notice that the same healing conditions were carried out for all the

interfacial systems, i.e. heating at 90°C for 1 hour, cooling down slowly to room

temperature, and keeping at room temperature for 24 hours before micro-droplet

debonding test. However, for some epoxy matrix systems, the glass transition

temperatures are higher than 90°C, i.e. the epoxy remains in a glassy state (Part II,

Fig. II-15). Some readers may doubt about why self-healing can occur even though

matrix is in solid state. In fact, this type of healing route relates mainly to the

exchange and the reforming of D-A covalent bonds which have an equilibrium

state at 90°C. Even though the debonded samples are heated above to 90°C, the

healing efficiency will not be improved due to: 1) as the healing temperature is


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higher than Tg, the chain segments of the epoxy matrix get a little mobility, which is

favorable for reforming D-A bonds. However, unlike hydrogen bonds or other

self-assembled bonds, D-A bonds formed by furan and maleimide groups tend to

separate at high temperature (>90°C). For this reason, there is no additional furan

groups that migrate near fiber surface and thus heating above 90°C is not required

for forming more D-A bonds; 2) Since the epoxy matrix is a crosslinked polymer,

the mobility of chain segments is limited, even if the material is heated up to the Tg

of matrix. However, if an appropriate pressure can be applied above Tg, a close

contact between the fiber surface and the epoxy matrix could be achieved, which

will be favorable for the formation of new D-A bonds. Nevertheless, for such

sample geometry (micro-droplet), the application of pressure is impossible. For

real composite materials, appropriate pressure could be applied during the healing

steps to improve the healing efficiency.

B.4 Conclusion

We have mentioned that the interfacial strength of D-A modified interphases is

mainly determined from the content of D-A covalent bonds formed between fiber

surface and the matrix. With the increasing of FGE content in the epoxy matrix, the

content of D-A covalent bonds increases gradually until the maleimide groups on

fiber surface are fully consumed. Above this value, i.e. about 40 wt.%, the

interfacial strength reaches a constant value. Furthermore, increasing FGE

content also leads to the decrease of the matrix crosslink density as the FGE

monomer acts as a chain extender between crosslinks. Of course, the crosslink

density of the interfacial layer also decreases with FGE content leading the

interphase to be tougher. The interfacial self-healing efficiency was calculated from


174

the interfacial shear strength (IFSS) and the critical energy release rate (Gic),

respectively. The results indicated that the healing efficiency exhibits an opposite

dependence with the initial interfacial strength, i.e. healing efficiency decreases

with the content of FGE. This phenomenon can be attributed to the cohesive

fracture process occurring at debonding. By considering both the interfacial

strength and the healing efficiency, we proposed that the optimum FGE content

should be from 20% to 30%. With such FGE content, a relatively strong

carbon/epoxy interphase with satisfactory interfacial self-healing ability can be built.

Additionally, the mechanical properties of epoxy matrix with such FGE content

can’t be weakened significantly, which confirms that such type of self-healable

interphase can be applied to (macroscopic) composite materials.


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General conclusions and perspectives

General conclusions:

This work was focused on the design of a self-healable carbon/epoxy interphase

with a multiple healing ability. For such objectives, thermally reversible Diels-Alder

covalent bonds were introduced between carbon fiber and epoxy matrix. This

manuscript is divided in three parts, including i) the choice of a healing strategy, ii)

the formation and characterization of interphase, and iii) the evaluation of

interfacial mechanical properties and healing ability.

Part I summarizes the background of this research, including the different

interfacial modification routes in composite materials, the characterization of

interfacial mechanical properties, and the self-healing techniques for polymer

material. After comparing the existing routes, the thermally reversible Diels-Alder

adducts formed between furan and maleimide groups were chosen to be

introduced into carbon/epoxy interphase as a healing incentive. For accurately

estimating interfacial self-healing ability, the micro-droplet debonding test was

chosen to characterize the mechanical properties of the interphase.

Part II describes the design of the self-healable interphase proposed in Part I.

The carbon fiber surface was treated using a three step protocol, HNO3 oxidization,

TEPA amination, and BMI grafting to get a maleimide-rich surface. The physical

and chemical changes of the fiber after each treatment step were characterized by

AFM, ATR-FTIR, XPS, and single fiber tensile test. We have concluded that

carbon fiber surface oxidization in nitric acid is an equilibrium process between

corrosion effect and oxidization effect. The oxidization effect induces an increase


176

of the content of activated functional groups on fiber surface depending on the

increasing of oxidization time. However, the corrosion effect erodes the fiber

surface leading to changes of fiber surface roughness. The fiber surface

roughness increases linearly within 60 minutes and decreases after 90 minutes, at

which some interconnected pores or internal voids and cracks are created in the

sub-surface. We have also shown that, within 90 minutes, oxidization treatment

does not damage fiber mechanical properties significantly. TEPA was confirmed to

be an appropriate amination agent allowing to introduce amino groups on fiber

surface, able to react with BMI. The results of ATR-FTIR characterization showed

that furan groups can be introduced in epoxy network by simply mixing

furan-epoxy monomer (FGE) into DGEBA prepolymer. The mass ratio of FGE can

impact the mechanical properties of epoxy and thus leading to a brittle matrix. By

combining the modified carbon fiber surface and matrix, a thermally reversible

self-healable interphase was built. After studying the reaction kinetics of D-A and

retro-D-A reactions by NMR, DSC ,and TGA, the best healing conditions were

determined as heating at 90 °C for one hour and cooling down slowly in furnace

until to the room temperature. The process of the micro-droplet debonding test was

also described in detail in this part.

To evaluate the impact of the D-A modification, the mechanical

characterization of some conventional carbon fiber/epoxy interphases (non-D-A

modified) was firstly carried out in Part III as reference systems. Some conclusions

are obtained. i) About sample shape parameters: the two contact angles (true (θ)

and apparent (ϕ)), formed between fiber and cured matrix droplet, have different

meanings for micro-droplet debonding test. The true contact angle (θ) can be used

to quantitatively compare the wettability of epoxy matrix on different carbon fiber


177

surfaces. However, the apparent contact angle (ϕ) is just a dimensional parameter

which influences the testing results because of the change of droplet profile. The

three dimensional parameters of matrix droplet such as embedded length (Le),

droplet diameter (Dd), and apparent contact angle (ϕ) determined by the volume of

matrix droplet have a linear relationship with each other. The distance between

droplet and paper support (free length, Lf) was proposed to be shorter than 250µm

to obtain a higher rate of test with success. The fiber diameter (Df) is not a constant

value (7µm) but distributed between 5µm and 10µm which influences the

calculating results. As a consequence, the actually measured values were used for

further calculation. ii) About interfacial mechanical properties: during the

micro-droplet debonding test, there are four rupture modes, including fiber

breakage, full interfacial completely debonding, cohesive debonding and matrix

fracture. The results obtained from interfacial completely debonding and cohesive

debonding can be used for analyzing the interfacial mechanical properties. All the

samples have undergone two debonding-healing cycles. Both the first and the

second interfacial debonding forces increase as a function of embedded length

with an obvious scattering. The interfacial shear strength (IFSS) measured by

droplet debonding test is not a constant but decreases with the increasing of

embedded length. iii) About interfacial self-healing ability: the conventional

carbon/epoxy interphase also has the ability to recover more or less interfacial

shear strength and the recovery ability is related to the interfacial structures.

However, this recovery is mainly attributed to the passive frictional force and thus

can’t be considered as an interfacial self-healing.

In Part IV, the properties of D-A modified interphases were studied. The

experiments were carried as a function of the fiber surface treatment and of the


178

network composition. The fiber-controlled experiments involve in combining epoxy

matrix with a serial of BMI-grafted carbon fibers which have been oxidized in nitric

acid for different times before grafting. The research results indicated that grafting

BMI on fiber surface worsened the wettability of epoxy matrix on fiber surface.

However, the interfacial bonding strength is increased because the D-A covalent

bonds are formed between fiber surface and matrix. The interphase composed

with T700-BMI which has been oxidized for 60 minutes during the first treatment

step possessed not only the highest interfacial bonding strength (89 MPa) but also

the highest healing efficiency (75%). The D-A modified interphase can be healed

several times and the healing efficiency is reduced step-by-step as the

debonding-healing cycle number is increasing. The matrix-controlled experiments

consist in comparing a serial of interphases combining carbon fiber T700-BMI

(60min) (carbon fiber oxidized for 60 minutes before grafting) and epoxy matrix

containing different mass ratio of FGE. The results confirmed that, as the

increasing of FGE content, the interfacial shear strength increased gradually until

to a constant value, for which the mass ratio of FGE was about 40% and the

maleimide groups on fiber surface were fully reacted to form D-A covalent bonds.

Furthermore, the increase of FGE content could also change the network density

of interfacial layer leading to a more flexible interphase. The interfacial self-healing

efficiency, which has been calculated with interfacial shear strength (IFSS) and

critical energy release rate (Gic), respectively, decreased with the increasing of

FGE content. By considering the bonding strength and healing efficiency at the

same time, the best mass ratio of FGE was confirmed to be between 20% and

30%.

Overall, this thesis is devoted to design a thermally self-healable


179

carbon/epoxy interphase based on thermally reversible Diels-Alder covalent bonds.

The designed interphase has not only the multiple self-healable abilities but also

the satisfactory mechanical properties. Additionally, the overall mechanical

properties of the composite material based on this interphase will not be weakened

significantly.

______________________________________________________________

Perspectives:

As the quantity of initial reactive groups for further grafting on the carbon fiber

surface is a key parameter, future works will focus on new fiber surface

treatment methods to graft more maleimide groups. In fact, the increase of the D-A

adducts quantity in interphase is an important issue to reach a higher healing

efficiency.

In addition, the designed interphase in this work is extended only on a

micrometric scale. Its properties, in comparison with a real composite material, are

very sensitive to the performance of components (fiber and matrix) and external

factors (environment, temperature, and humidity etc.). So, building a real

carbon/epoxy composite material with self-healable interphase is necessary. Here,

we have to indicate that the micro-droplet debonding test perfectly relevant to

characterize the mechanical performance of the designed interphase involves the

entirely breaking of the interphase. However, in a real composite, the concept of

interfacial self-healing is generally applied to extend the service life and thus

should take place at the initial stage of crack. In this case, the damage of

interphase is generally not so serious, which should be easier to be healed than

the damage of micro-droplet debonding test. For this reason, some other more


180

artful testing methods should be proposed.

Furthermore, transferring the same concept to other composite systems

such as carbon nanotube based nano-composite materials will be also very

interesting. Although the designed interphase already shows a favorable

self-healing ability, it needs to introduce mono-functional monomers into epoxy

network which will weaken the mechanical properties of matrix and also the

overall mechanical properties of composite material. For this reason, a new

interfacial self-healing system without weakening the performance of matrix and

fibers would be established.

In fact, interphase is only one of the key points for a composite material.

Other essential factors such as fiber and matrix are also very important and tend to

fracture during application. If we can design a composite material and make all the

components have a self-healing ability, it will be very effective and useful for many

structural composite part in a lot of industrial fields like transport (automotive,

railway, aerospace…), sport, energy, …


181

Annexe

Annex I. Interfacial failure criteria

Local interfacial shear strength (τd)

𝐹 =2𝜋𝑟𝑓𝛽

⎩⎪⎨

⎪⎧ 𝜏𝑑𝑡𝑡𝑡ℎ[𝛽(𝑙𝑒 − 𝑡)] −

𝜏𝑇𝑡𝑡𝑡ℎ[𝛽(𝑙𝑒 − 𝑡)] ∗ 𝑡𝑡𝑡ℎ �[𝛽(𝑙𝑒 − 𝑡)]

2�

+𝛽𝑡𝜏𝑓 ⎭⎪⎬

⎪⎫

𝛽2 =2

𝑟𝑓2𝐸𝐴𝐸𝑚⎝

⎛ 𝐸𝐴𝑉𝑓 + 𝐸𝑚𝑉𝑚𝑉𝑚

4𝐺𝐴+ 1

2𝐺𝑚� 1𝑉𝑚

𝑙𝑡 1𝑉𝑓− 1 −

𝑉𝑓2 �⎠

⎞

𝜏𝑇 = 𝐸𝐴𝛽𝑟𝑓(𝛼𝐴 − 𝛼𝑚)∆𝑇/2

τf =f

2πrr𝑡

F — Current shear force f — Current frictional force

GA — Fiber shear moduli Gm — Matrix shear moduli

le — Embedded length a — Current crack length

rf — Fiber radius β — Nayfeh’s shear-lag parameter

τT — Thermal shrinkage stress τf — Frictional stress

EA — Fiber axial tensile moduli Em — Matrix tensile moduli

Vf — Fiber volume fraction Vm — Matrix volume fraction


182

αA — Fiber axial thermal expansion coefficient

αm — Matrix thermal expansion coeffcient

ΔT — Difference between testing temperature and stress-free temperature

Critical energy release rate (Gic) and adhesion pressure (σd)

Simple analysis

𝑮𝒆𝒆 =𝑟𝑓2�𝐶33𝑠𝜎𝑑2 + 2𝐷3𝑠𝜎𝑑∆𝑇 +

𝐷32

𝐶33∆𝑇2�

Shear-lag model

𝝈𝒅 =𝐹𝑑

2𝜋𝑟𝑓2= −

𝐷3𝑠𝛥𝑇𝐶33𝑠

+ �2𝐺𝑖𝑖𝑟𝑓𝐶33𝑠

𝑪𝟑𝟑𝟑 = 12

( 1𝐸𝐹

+ 𝑉𝑓𝑉𝑚𝐸𝑚

) 𝑪𝟑𝟑 = 𝟏𝟐� 𝟏𝑬𝑭

+ 𝑽𝒇𝑽𝒎𝑬𝒎

� − 𝑽𝒎𝑨𝟑𝟐

𝑽𝒇𝑨𝟎

𝑫𝟑𝟑 = 𝟏𝟐

(𝜶𝒇 − 𝜶𝒎) 𝑫𝟑 = −𝑽𝒎𝑨𝟑𝑽𝒇𝑨𝟎

(𝜶𝑻 − 𝜶𝒎) + 𝟏𝟐

(𝜶𝒇 − 𝜶𝒎)

𝑨𝟎 = 𝑽𝒎(𝟏−𝝊𝑻)𝑽𝒇𝑬𝑻

+ 𝟏−𝝊𝒎𝑬𝒎

+ 𝟏−𝝊𝒎𝑽𝒇𝑬𝒎

𝑨𝟑 = −(𝝊𝒇𝑬𝒇

+ 𝑽𝒇𝝊𝒎𝑽𝒎𝑬𝒎

)

𝑽𝒇 = 𝟑𝟐�𝒅𝒇𝑫�𝟐

𝑽𝒎 = 𝟏 − 𝑽𝒇

Vf — Fiber volume fraction Vm — Matrix volume fraction

EF — Fiber axial moduli ET — Fiber transverse moduli

Em — Matrix moduli

υf — Fiber axial Poisson’s ratio υT — Fiber transverse Poisson’s ratio


183

υm — Matrix Poisson’s ratio

D — Droplet diameter df — Fiber diameter

le — Droplet length (embedded length)

αf — Fiber axial thermal expansion coefficient

αT — Fiber transverse thermal expansion coefficient

αm — Matrix thermal expansion coefficient

ΔT — Difference between testing temperature and stress-free temperature

Annex II. Characterization techniques for carbon fiber surface

and matrix

Atomic Force Microscopy (AFM)

AFM surface patterns were acquired in air at room temperature using a

Nanoscope IIIa Multimode (Digital Instruments/Brucker, CA). Tapping mode was

performed at a scan rate of 1 Hz with uncoated silicon probes having a resonance

frequency included between 280 and 405 kHz, a spring constant between 20 and

80 N/m, a length between 115 and 135 µm, width between 30 and 40 µm. The

surface topography images and the average roughness were analyzed by

Nanoscope 6.14R1 software.

Attenuated total reflectance Fourrier transform infrared spectroscopy

(ATR-FTIR)

ATR-FTIR is relevant to characterize the functional groups created after the

surface treatments. The test sample was fiber bundles that was compacted and


184

light-tight. The spectra were recorded bythe Nicolet iS10 Thermo Scientific

spectrometer with a ZnSe crystal (Ge critstal for matrix characterization) at room

temperature. The resolution is 4cm-1 and 32 scans were collected for each sample.

X-ray photoelectron spectroscopy (XPS)

The XPS spectra were collected on an AXIS ULTRADLD (KRATOS ANALYTICAL)

X-ray photoelectron spectrometer. Data were collected using Slot (300X700μm)

mode with the different pass energy of 160eV (full range) and 20eV (region). The

X-ray source was Al monochromatic (1486.6 eV, 150W) and the energy scale of

the spectrometer was 0-1200eV. The C1s binding energy (BE) of the graphitic

peak was fixed at 284.6eV for calibration.

Monofilament tensile test

The monofilament tensile test was carried out on a tensile testing machine using a

10N maximum load cell with a sensitivity of 0.001N. The displacement rate and the

Data collection rate were 0.1mm/min and 100Hz respectively. Since the carbon

monofilament is very thin and fragile, it’s difficult to clamp on testing machine

directly. Thus a specific sample, shown as following, was used. The two ends of

monofilament were glued on a paper frame with a square hole in center. When

starting a test, the upper and lower ends of the paper frame will be clamped with

jigs, by which the monofilament can be effectively supported. Then the left and

right sides will be cut off carefully.


185

Nuclear magnetic resonance (NMR)

The 1H NMR spectra was recorded on a Bruker AVANCE250 (250MHz) or

AVANCE400 (400MHz) spectrometer using DMSO-d6.


186


187

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