Thermo-hydrolytic resistance of polyepoxide–glass fibres interfaces by the microbond test

Composites Science and Technology 68 (2008) 2028–2033

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Thermo-hydrolytic resistance of polyepoxide–glass fibres interfaces by the microbond test Philippe Zinck a,*, Jean-Francßois Gérard b a

Synthèses Organométalliques et Catalyse, Unité de Catalyse et Chimie du Solide, UMR CNRS 8181, ENSCL, Cité Scientifique, 59652 Villeneuve d’Ascq, France Laboratoire des Matériaux Macromoléculaires, Ingénierie des Matériaux Polymères, IMP, UMR CNRS 5627, Université de Lyon, INSA de Lyon, Bât Jules Verne, 20 Av. A. Einstein, 69621 Villeurbanne Cedex, France

b

a r t i c l e

i n f o

Article history: Received 25 July 2007 Received in revised form 12 February 2008 Accepted 23 February 2008 Available online 29 February 2008

Keywords: A. Polymer–matrix composites B. Fibre–matrix bond B. Interphase B. Hygrothermal effect B. Durability

a b s t r a c t The use of the microbond test as a tool for the characterization of the thermo-hydrolytic resistance of interfacial zones is critically discussed for polyepoxide based composites materials. It is shown that the differences of properties observed between macroscopic and microscopic scales can induce a difference in aging conditions. The energy release rate vs. time of exposure at 60 °C and 98% RH is reported for epoxyde–amine and epoxyde–anhydride/glass fiber systems. The evolution of the energy release rate is interpreted in terms of different degradation mechanisms: release of internal stresses and plasticization for the short term behavior of both systems, while interfacial hydrolysis is advanced for the long term behavior of the epoxyde–anhydride/glass fiber system. This latter chemical degradation is attributed to the presence of hydrolysis sensitive functional groups in the interphase resulting from the reaction between the anhydride co-monomer and the organosilane coupling agent. Trends for the design of thermo-hydrolytic resistant interphase are given on this basis. Ó 2008 Elsevier Ltd. All rights reserved.

1. Introduction Owing to their excellent mechanical properties, composites materials have been widely used all along the last four decades. Their use in hostile environment has given rise to studies devoted to the durability of their properties. In fact, as those materials are subjected to high temperature and moisture/water attack, different mechanisms of degradation occur simultaneously according to the severity of the exposure conditions: - Mechanical degradation, including plasticization of the matrix and swelling, release of internal stresses [1,2]. - Irreversible chemical degradation, such as hydrolysis of the matrix and the interphase [3–5]. The prediction of the long term behavior of these materials still remains difficult, due to the lack of knowledge of the long term behavior of the interface/interphase. The study of such zones is indeed difficult due to the small scale involved. Dynamical mechanical analyses were carried out to study the thermo-hydrolytic resistance of glass fiber sizing in situ [6]. Macroscopic samples representative of the interfacial zones were synthesized in order to * Corresponding author. E-mail addresses: [email protected] (P. Zinck), [email protected] (J.-F. Gérard). 0266-3538/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.compscitech.2008.02.025

characterize their thermo-hydrolytic resistance [7,8]. The characterization of single interphase was reported in most cases through a mechanical response. The evolution of the interfacial load-transfer capacity in an aggressive environment was mainly studied through micromechanical testing such as pull-out [9,10], microbond test [11–17], single fiber fragmentation test [18–23] and nanoindentation [24–26]. It should be noted here that, from a practical point of view, micromechanical testing using microcomposite provides a less time consuming method for the characterization of hydrothermal aging than studies performed at the macroscopic scale. Indeed, due to the small scale involved, aging studies can be carried out in several days vs. several months for macroscopic samples. The extrapolation of microscopic results to macroscopic behavior has in turn to be considered with caution in the case of the microbond test. In fact, polymeric microdroplets can exhibit properties different from that of the bulk matrix. The specific geometry of the specimens can lead to a loss of stoechiometry, especially in the case of polyepoxide matrices [27,28]. Losses in glass transition temperature were reported for polyepoxide microdroplets based on a diglycidyl ether of bisphenol A (DGEBA) epoxyde prepolymer combined to aliphatic and aromatic amines co-monomer as well as anhydrides. It is well known that the stoichiometric ratio plays an important role in water absorption processes and a lack of stoichiometry can lead to drastic modification of the sorption behavior and degradation mechanisms of polyepoxide networks [29–31]. It is likely in this frame that aging

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and determined at the onset of the transition zone. Young’s modulus and Poisson‘s ratio used for modeling the experimental data were determined in tension on a Adamel Lhomarghy DY22 apparatus at a cross-speed of 0.5 mm/min. The yield point was measured in compression using the same apparatus. The thermal expansion coefficient were determined on a DMA 2980 (TA Instrument) using a penetration clamp at 2 °C/min. Physical and mechanical properties of the networks used for the calculation of the energy release rater are given in Appendix A. Details concerning synthesis of model films, preparation of microbond specimens and shearing device can be found in previous studies [28,32]. Hydrothermal aging experiments were carried out by exposing model films and microcomposites to 98% RH at 60 °C for different periods. Microbond specimens were tested at a cross-speed of 0.5 mm/min, and data were fitted by the model of Scheer and Nairn [33] (the formalism is presented in Appendix A). A minimum of 20 debonded specimens were considered for the calculation of the energy release rate. A typical maximum force vs. embedded area plot is represented in Fig. 1.

conditions of microcomposites differ from that of the macroscopic level, and in turn, the environment and degradation of the interface. The evaluation of such differences has never been reported so far to our knowledge, and will be discussed in this work. A comparison between the plasticization of microscale specimens vs. bulk specimens is proposed using thin films, models of the microcomposites. The ability of the microbond test to characterize the thermo-hydrolytic resistance of interfaces in composites materials is reported for two polyepoxides/glass fibres systems. Factors governing the thermo-hydrolytic resistance of polyepoxide/glass fiber interphase are finally discussed. 2. Experimental section E-glass fibres have been supplied by VETROTEX Int. The sizing referred to as P122 by Vetrotex Co. is known as a universal sizing suitable for polyepoxide as well as polyester matrices. The coupling agent is the c-aminopropyltriethoxysilane. The main properties of the fibres are given in Appendix A. The two types of polyepoxide networks studied are based on DGEBA (diglycidyl ether of bisphenol A, Ciba Geigy, n = 0.15). Two different co-monomers were selected, 4,40 -methylenebis[2,6-diethyleneaniline] (MDEA, Lonza) and anhydride cis-4methyl-1,2,3,6-tetrahydrophthallic (MTHPA, Anchor Chemicals). The chain polymerization between epoxyde and anhydride is more convenient in the presence of a Lewis base. A tertiary amine was selected for this purpose, 2,4,6 (dimethyl aminomethylene) phenol (DMP30, Fluka). Chemical formulae are presented Table 1. All reactions were conducted for radii epoxyde/anhydride = 1 and epoxyde/amine = 1. The amount of Lewis base used in the epoxyde/ anhydride reaction was set after study to 1,5% wt. The chosen radii and cure schedule (4 h 135 °C, 4 h 190 °C for DGEBA/MDEA and 1 h 100 °C, 5 h 160 °C for DGEBA/MTHPA, heating rate of 1 °C/min and cooling rate of 2 °C/min) lead to fully cross-linked networks as revealed by differential scanning calorimetry thermograms. Glass transition temperature were measured by differential scanning calorimetry at 10 °C/min under nitrogen atmosphere

Maximal force (N)

0.6

0.4

0.2

0.0 0

2000

4000

6000

8000

10000

Embedded area (µm2) Fig. 1. Maximal force vs. embedded area for DGEBA–MDEA/glass fiber microdroplets after 1 h in an oven at 60 °C and 98% HR.

Table 1 Co-monomers and initiator used for the synthesis of polyepoxide networks Name

Representation

DGEBA n = 0.15

CH2 CH CH2 O O MTHPA

CH3

CH3

CH3

O CH2 CH CH2 O

C

OH

CH3

CO CO CO

DMP30

CH3 CH3

OH CH2

N CH2

CH2 MDEA

C2H5 N2 H C2H5

N

N

CH3 CH3

CH3 CH3 C2H5

CH2

NH2 C2H5

n

C CH3

O CH2 CH CH2 O

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3. Results and discussion 3.1. Water induced plasticization: microscopic vs. macroscopic behavior of polyepoxide networks Films were used as models of polyepoxide microdroplets in order to show and tentatively quantify differences between microscopic and macroscopic properties in terms of stoichiometry, physical and mechanical properties [28]. The glass transition temperatures of the films were wound to be ca. 50 °C lower than those of the bulk matrices. From this point, we became aware of the possibility of observing aging conditions that are more drastic for microscale specimens vs. macroscopic specimens, since one of the key parameter of hydrothermal aging is the difference between the glass transition temperature and the aging temperature. We measured in this frame the glass transition temperatures of the model films after a thermo-hydrolytic treatment in the same experimental conditions as those used for the exposure of microdroplets (Section 3.2). The glass transition temperature of macroscopic specimens at the initial state and at equilibrium after 30h1/2 in distillated water at 60 °C [34] are also represented Fig. 2 for convenience. The following statements can be made: (i) The difference between the glass transition temperature and the aging temperature is more important for the films than for bulk specimens, particularly for the amine based polyepoxide network, highlighting aging conditions that are more severe as compared to macroscopic samples. (ii) The loss in glass transition temperature after water sorption is more important in the microscopic configuration than in the macroscopic one, highlighting degradation mechanisms of higher amplitude on the microscale. (iii) The glass transition temperature of aged films are lower than the aging temperature; the polymeric film is thus at the rubbery state during the thermo-hydrolytic treatment when in fact macroscopic samples are at the glassy state during exposure at the same temperature.

3.2. Interphase thermo-hydrolytic resistance and mechanisms of degradation Fig. 3 presents the evolution of the energy release rate of glass fiber/polyepoxide interfaces derived from microbond experiments with respect to time of exposure at 60 °C and 98% HR. For both networks, a loss in fracture energy is observed after one hour exposure, followed by a plateau for the epoxyde–amine based system. A long term degradation is additionally observed for the epoxyde–anhydride system before the plateau. Such evolutions are in agreement with results of the literature [9,10,12,15]. The initial loss in fracture energy is attributed to the release of internal stresses together with the plasticization of the matrix, while long term degradations are attributed to a chemical degradation of interfacial zones. Concerning time scales, the time required to reach saturation for a microdroplet at 60 °C and 98% RH can be roughly estimated using the diffusivities determined for the bulk matrices a the same temperature [34]. This leads to a value of a few minutes. Table 2 Calculated amount of water (wt.%) absorbed by model films (film 24 h, film 48 h) compared to experimental values measured on macroscopic specimens (bulk, after [34]) System

Bulk

Film 24 h

Film 48 h

MTHPA MDEA

1.0 1.8

2.3 7.5

2.2 7.4

180 MTHPA

160 Aging Temp.

120 100

Room Temp.

80 60 40 20 0

250

MDEA

140

Bulk dry

Bulk eq.

Film dry

Film 24h

Film 48h

Fig. 2. Glass transition temperatures of model films vs. bulk specimens at the dry state (bulk dry and film dry, respectively), at the equilibrium after sorption at 60 °C in distillated water for macroscopic specimens (bulk eq.) and after 24 and 48 h exposure at 60 °C, 98% HR for the model films (film 24 h and Film 48 h, respectively).

Fracture energy (J/m2)

Glass transition temperature (˚C)

The higher loss in glass transition temperature is probably due to a higher water uptake at equilibrium and resulting plasticization. Hydrolysis could indeed not be detected on bulk samples [34], and the stability of the glass transition temperature of the films between 24 h and 48 h acts also for the absence of chemical degradation of the films. We further use polymer–diluents laws to evaluate the quantity of water which would have been absorbed by

the films under the assumption that chain scission does not occur. Thermo-gravimetric analyses were indeed too difficult to realize due to the small scale involved. Fox’s law of mixtures [35], and other as approaches based on free volume considerations of Kelley and Bueche [37–39] were compared, and two models were selected on the basis of the quality of their prediction of the glass temperature transition for the bulk samples (more details can be found in the Appendix A). The amount of water absorbed by the films was estimated from the glass transition temperature using these models. Results are given Table 2 in comparison with the bulk. Anhydride based thin films could absorb twice much water than bulk matrices, and amine based films four times more. These observations show that the plasticization and aging conditions of polyepoxide matrices at the microscale can be more severe, and possibly increase the extent of degradation of interfacial zones in microdroplets as a consequence. This should be considered with caution however, since the relevance of thin films as model of microbond specimens is difficult to assess, in particular in the frame of thermo-hydrolytic studies.

DGEBA/MTHPA 200 150 100 DGEBA/MDEA

50 0 0

30

60

90

120

Exposure time (h) Fig. 3. Energy release rate for DGEBA–MTHPA (triangles, black) and DGEBA–MDEA (circles, grey)/glass fibers microdroplets as a function of the time of exposure at 60 °C and 98% HR.

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It is thus reasonable to consider that the equilibrium is reached after 1 h exposure. The initial loss is more important for amine based microcomposites. Both release of residual stresses and plasticization of the network can explain this difference. The loss in glass transition temperature and the subsequent weight gain at equilibrium show that the extent of plasticization is higher for epoxyde–amine systems. Thermal residual stresses are a function of Young modulus, thermal expansion coefficient of the matrix at the glassy state and the difference between the glass transition temperature and the room temperature. As detailed in Appendix A, a comparative evaluation of residual stresses can be performed using the one dimensional model from Tsai & Hahn [40]. The calculated thermal residual stresses are more important by a factor 2 for amine based microcomposites vs. anhydride based microcomposites, around 50

and 25 MPa, respectively. This trend has been confirmed by photoelastic studies on model DGEBA/MDEA and DGEBA/MTHPA/glass inclusions systems [41]. If the epoxyde–amine/glass fiber interphase does not undergo any additional degradation, we attribute the gradual loss in energy release rate observed at longer times for anhydride based microcomposites to a chemical degradation of the interphase, as proposed in the literature, and keeping in mind that both polyepoxide networks do not undergo any chemical degradation in the same experimental conditions [34]. The hydrolysis of specific functional groups formed in epoxyde–anhydride/glass fiber interphase may be advanced, as detailed hereafter. The structure of interphase results from miscibility/diffusion of the co-monomers with/in the sizing of glass fibers and subsequent chemical reactions [42]. The main component of a sizing is the cou-

Table 3 Functional groups in epoxyde–amine and epoxyde–anhydride polyepoxide networks (bulk) and in the interphase between polyepoxides and aminosilane on a substrate Bulk

Epoxyde Amine

CH

CH2

OH

+

N

NH2 OH

O

CH Epoxyde Anhydride

CH2

O

O

+

CH

O

CH2

OCO

OCO O

Interphase

Epoxyde AmineOSil

CH

CH2

OH

+

OH

O

Amine AmineOSil

Anhydride AmineOSil

NH2

+

N

NH2

CH

CH2

O

H

C

N

NH2

O

O

+

NH2

O O N

O The organosilane amine group is referred to as AmineOSil.

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pling agent, and in the case of industrial sizing, it is assumed that silane coupling agent migrates to the interface providing an interfacial region which is similar to that obtained starting from pure coupling agent solutions [43]. A three-dimensional graded crosslinked network is representative of the structure in both cases [6,42,43]. The coupling agent present in the P122 sizing is the c-aminopropylthriethoxysilane, and amide and imide functional groups [44] have been observed in interphase between epoxyde– anhydride network and c-aminopropylthriethoxysilane on a substrate. These groups result from the reaction between the anhydride co-monomer and the amine group of the silane, and are not present in epoxyde–amine/glass fiber interphase. The sensitivity of amide and imide groups to hydrolysis is well known, particularly in the case of oligomers/polymers, which results in significant decrease of the mechanical properties [45,46]. We proposed to attribute the behavior observed here to the hydrolysis of such functional groups present in epoxyde–anhydride/glass fiber interphase. Functional groups present in epoxyde–amine/glass fiber interphase are in turn similar to those present in the bulk, since the organosilane and one of the co-monomers bear the same chemical function, an amine. Such considerations are tentatively summarized Table 3. Readers interested in epoxyde - aromatic amine polycondensation and epoxyde - anhydride chain polymerization are invited to consult Refs. [47,48], respectively, for more details. 3.3. Design of thermo-hydrolytic resistant interphase and composites materials We should finally address the applicative issue of thermohydrolytic resistant interphase and composite materials design. Anhydride based polyepoxide networks are valuable starting materials for use in water environment, as they exhibit smaller water uptake and weaker interactions of specific functional groups with water molecules as compared to epoxyde–aromatic amine networks [34]. Their use as matrices for high performance composites materials is thus promising. This should however be done in combination with fiber sizing that do not contain coupling agents bearing a functional group able to react with the anhydride co-monomer. The generation of hydrolysis sensitive groups in interfacial zones is indeed detrimental, as highlighted by the diglycidyl ether of bisphenol A/anhydride cis-4-methyl-1,2,3,6-tetrahydrophthallic/c-aminopropylthriethoxysilane combination. 4. Conclusion The use of the microbond test as a tool for the characterization of the thermo-hydrolytic resistance of interfacial zones has been evaluated for polyepoxide based composites materials. It is shown using model films that water induced glass transition temperature loss can be more important on a microscopic scale. Such differences between macroscopic and microscopic properties could possibly induce more severe aging conditions for microcomposite specimens. The evolution of the energy release rate vs. time of exposure at 60 °C and 98% RH is reported for epoxyde–amine and epoxyde– anhydride/glass fiber systems. Different degradation mechanisms are advanced in order to explain the observed trends: release of internal stresses and plasticization for the short term degradations of both systems, and interfacial hydrolysis for the long term behavior of the epoxyde–anhydride system. This chemical degradation is proposed to be a consequence of the presence of hydrolysis sensitive functional groups such as amide or imide in the interphase resulting from the reaction between the anhydride co-monomer and the aminosilane coupling agent. The issue of thermo-hydrolytic resistant interphase and composites materials design is discussed in this frame. The use of epoxyde–anhydride

matrices for high performance composites materials is promising in combination with a fiber sizing that do not contain coupling agents bearing a functional group able to react with the anhydride co-monomer. It is finally noteworthy that, in spite of the possible more severe aging conditions at the microscale in the case of polyepoxide, the evolution of the properties of interfacial zones in the course of a thermo-hydrolytic treatment can effectively be addressed via the microbond test. Acknowledgements DER/EDF is acknowledged for its financial support and Vetrotex International for supplying E-glass fibers. Appendix A A.1. Materials properties E-glass fibres: diameter 19.1 ± 1.4 lm; tensile modulus 73 GPa; Poisson coefficient 0.22; thermal expansion factor 5.106 K1. Polyepoxides: glass transition temperature Tg, Young modulus E, Poisson coefficient m, yield point ry, and thermal expansion coefficients at the vitreous and rubbery state, av and ac, respectively. Matrix

Tg (°C)

E (GPa)

m

ry (MPa)

av  104 (K1)

ac  104 (K1)

DGEBA/MDEA DGEBA/MTHPA

158 124

3.35 3.89

0.41 0.37

105.5 108.4

1.18 0.75

2.19 2.25

A.2. Model of Scheer and Nairn [33] The simplified analysis derived from variational mechanics leads to vffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ! u D3s DT u 2Gic DT 2 D23s D23 t þ þ  rd ðqÞ ¼  C 33s rf C 33s C 33s C 33s C 33 where rd is the stress at decohesion and q the axial ration (embedded length/fiber diameter), the C and D constants depends on the sample dimensions and on the mechanical properties of the fiber and the matrix:   1 1 Vf C 33s ¼ þ 2 Ef V m Em D3s ¼

1 ðaf  am Þ 2

C 33 ¼

  1 1 Vf V m A23 þ  2 Ef V m Em V f A0

D3 ¼ 

A0 ¼

V m A3 1 ½aT  am  þ ðaf  am Þ 2 V f A0

V m ð1  mT Þ 1  mm 1 þ mm þ þ V f ET Em V f Em

  mf V f mm A3 ¼  þ Ef V m Em In these equations, Ef and ET are the axial and transverse tensile moduli of the fibre, mf and mT the axial and transverse Poisson’s ratio of the fibre, af and aT the axial and transverse thermal expansion coefficients of the fibre, Em, mm and am the tensile modulus, Poisson’s ratio and thermal expansion coefficient of the matrix.

P. Zinck, J.-F. Gérard / Composites Science and Technology 68 (2008) 2028–2033

Vf and Vm are the volume fraction of the fiber and the matrix, respectively. According to the works of Scheer and Nairn, they were calculated from the plot measured droplet diameter vs. measured droplet length. A.3. Models describing the plasticization of a polymer A.3.1. Model of Fox [35] 1 w1 w2 ¼ þ T g T g1 T g2

A.3.2. Model of Kelley and Bueche [37] ap V p T gp þ ad ð1  V p ÞT gd ap V p þ ad ð1  V p Þ

a ¼ al  av A.3.3. Model of Morgan and O’Neal [38]

A.3.4. of Mc Kague et al. [39] avp V p T gp þ ald ð1  V p ÞT gd avp V p þ ald ð1  V p Þ

A.3.5. Numerical treatment The following values were taken in the calculation: adl = 5.23 104 K1 at 60 °C and avp = 0.68  104 K1 (water at 0 °C, from Ref. [51]) and Tgd = 137 °C [52]. Matrix

Tg exp.

Fox

Kelley and Bueche

Morgan and O’Neal

Mc Kague et al.

MTHPA MDEA

102 124

117 142

112 126

121 152

104 133

The approach of Kelley and Bueche and Mc Kague et al. were selected for the amine and anhydride systems, respectively. A.4. Evaluation of thermal residual stresses in a microbond specimen Models for built-in residual stresses in composites materials have been reviewed by Wagner [49]. The most reliable models were those of Hahn and Tsai [40] and Nairn [50]. It was pointed out that the presence of inhomogeneous cooling stresses in a polymeric droplet largely modifies the stress state and that this consideration has to be taken into account in the modeling. Since our attempt is only to compare two matrices, we chose the one dimensional approach of Tsai and Hahn for its simplicity. It is considered that the sole residual stresses components present in the fibers and in the matrix are the longitudinal one, given by ¼ ðam  af ÞðT  T ref Þ

Ef 1 þ ðuumf ÞðEEmf Þ

by assuming the microdrop to be an ellipsoid of revolution. We take for the numerical treatment typical value of 10 lm for the fiber radius rf and 50 lm for the microdroplet radius rm. This gives a fiber volume fraction of 6% (which is actually low compared to unidirectional composites).

[1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]

T g ¼ V p T gp þ ð1  V p ÞT gd

rfz

where a and E are the coefficient of thermal expansion and Young’s modulus, respectively, u is the content by volume, with the subscript m or f for the matrix or the fiber, T is the test temperature and Tref the stress free reference temperature. The volume fraction of the fiber in a microbond specimen were reported [33] to be

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where Tg, Tg1 and Tg2 are the glass transition temperature of the polymer–diluant system, the polymer and the diluant and w1 and w2 their respective weight fraction in the mixture.

Tg ¼

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