Carbon fiber-reinforced epoxy filament-wound composite laminates exposed to hygrothermal conditioning José Humberto S. Almeida Jr.1 • Samia D. B. Souza2 • Edson C. Botelho2 • Sandro C. Amico1 Received: 14 November 2015 / Accepted: 25 January 2016 / Published online: 5 February 2016 � Springer Science+Business Media New York 2016 Abstract This study focuses on the evaluation of the effect of hygrothermal conditioning on tensile, compres- sive, in-plane and interlaminar shear properties, and also on the viscoelastic characteristics of carbon fiber/epoxy lam- inates. Flat unidirectional laminates were manufactured by dry filament winding and cured under hot compression. The laminates were later exposed to hygrothermal condi- tioning in a chamber, following the recommendations of ASTM D5229M. All composite coupons were tested before and after conditioning. An analytical Fickian model was used to fit experimental data, showing very good estimates. Shear strength and modulus reduced to about 30 and 38 %, respectively. All specimens presented acceptable failure modes; shear specimens failed at the gage section with delaminations and fiber/matrix debonding, whereas short beam specimens failed via delaminations at the specimen mid-plane. Moisture penetration through the carbon/epoxy surface lead to interfacial debonding and matrix plasti- cization. Puck’s failure envelope accurately predicted failure under compressive and shear loading. Introduction Carbon fiber-reinforced epoxy composites are commonly used in applications with high structural demand, such as marine, aeronautical, and aerospace sectors [1]. Among the manufacturing processes available for advanced polymer composites, filament winding (FW) stands out, being able to produce from flat laminates to elbows and curved-sur- face structures. The high fiber volumetric fraction and precision in angle deposition, along with the use of a ten- sioned and continuous reinforcement are the main features of FW. The oil and gas sector, in particular, is increasingly replacing metallic-based structures with composites based on the need to increase payload in marine structures by reducing weight and increasing corrosion resistance. Indeed, pipelines, subsea systems, and submersible struc- tures operate under aggressive environments, which mod- ify the mechanical response of the composite, decreasing damage tolerance and sometimes leading to premature failure [2]. Prediction of failure and damage in composites is typi- cally complex mainly due to their orthotropic characteristic, and this becomes even more complex when the material is under extreme weathering conditions. When under these environments, composites can absorb moisture, which affects their long-term durability and properties [3]. Since high-performance carbon fibers absorb virtually no mois- ture, absorption is largely matrix-dominated [4], suggesting that matrix-dominated failure mechanisms will be more significantly affected. Also, if the matrix content is high or if a fiber more prone to absorption is being used (e.g., pol- yaramid), the greater the expected water absorption. Moisture permeation is dominated by diffusion, capil- larity, and/or transport by micro-cracks, and the rate of & José Humberto S. Almeida Jr. jhsajunior@globomail.com 1 PPGE3M, Federal University of Rio Grande do Sul, Av. Bento Gonçalves 9500, Porto Alegre, RS 91501-970, Brazil 2 Department of Materials and Technology, Universidade Estadual Paulista (UNESP), Av. Ariberto Pereira da Cunha 333, Guaratinguetá, SP 12516-410, Brazil 123 J Mater Sci (2016) 51:4697–4708 DOI 10.1007/s10853-016-9787-9 http://orcid.org/0000-0002-9408-7674 http://crossmark.crossref.org/dialog/?doi=10.1007/s10853-016-9787-9&domain=pdf http://crossmark.crossref.org/dialog/?doi=10.1007/s10853-016-9787-9&domain=pdf moisture uptake varies with the type of matrix, fiber ori- entation, water temperature, moisture [5], and fiber content (Vf). Moisture wicking can promote changes in mechanical, thermo-mechanical, and thermo-physical characteristics of the matrix by plasticization, swelling, cracking, and hydrolysis, and it can also degrade fiber-matrix bonding [3, 6]. When combined with high temperatures, it may induce irreversible structural changes in the polymer network. The extent of damage may change swelling, since cracking and blistering cause higher uptake, whereas the leaching of small molecule components results in gradual decrease in weight gain [7]. Ultraviolet radiation (UV) can also degrade the matrix and cause irreversible structural changes. UV radiation can trigger photolytic, photo-oxidative, and thermo-oxidative reactions within the matrix. Degradation can cause from simple discoloration to substantial loss of mechanical properties as a result of C–C polymer chain scission [8] due to photo-oxidation [9]. Another effect is the cross-linking, which restricts molecular mobility and reduces the ability of the material to accommodate externally applied strain [10]. Marine structures sometimes operate under combined UV and seawater influence both being detrimental to polymeric composites. Botelho et al. [11] evaluated the effect of different hygrothermal conditionings on the shear behavior of carbon/epoxy composites and Mouzakis et al. [9] studied accelerated environmental aging of glass/ polyester composites, reporting a significant decrease in compression and shear strength. Thus, this paper addresses the effect of environmental conditioning (accelerated water immersion) on the tensile, compressive, in-plane shear, and viscoelastic properties of flat filament-wound carbon fiber/ epoxy laminates. Experimental Materials and manufacturing Carbon fiber/epoxy prepreg tow (towpreg) from TCR Composites, with Toray T700-12 K-50C carbon fiber and UF3369 epoxy resin system, was used in this work. The flat laminates were manufactured using a rectangular stainless steel mandrel (327 9 228 9 12 mm3) and a KUKA robot KR 140 L100 with control and peripheral devices from MFTech, with an achievable tolerance in the winding angle lower than 0.5 %. Hoop flat laminates were produced on the designed mandrel, as presented in Fig. 1. Two simul- taneous rovings (towpregs) were used to produce the laminates. The towpregs were then wound onto the man- drel. After winding, a polyester-based shrink tape was used to wrap and help consolidating the laminate during the curing process that followed. The tape begins to shrink around 70 �C and reaches a maximum shrink force at 150 �C. This aids ply compaction and the elimination of voids, yielding a resin-rich part surface. The laminates were cured by hot compression under six ton for 4 h at 130 �C. The final fiber volume fraction was &72 % (measured by acid digestion following ASTM D3171-11) and mean ply thickness was 0.35 mm in the 4-layer laminate (overall thickness of 1.40 mm). Weathering The carbon fiber/epoxy specimens were exposed to a combination of temperature and humidity for 60 days at 80 �C under a relative humidity of 90 % inside a hygrothermal climatic chamber Marconi, model MA835/ UR. The hygrothermal exposure was carried out following the recommendations of ASTM D5229M-14. The samples went through a drying procedure to guarantee that they were dry and in mass equilibrium before aging, preventing that initial and atmosphere humidity affected water uptake values [1, 12]. The drying cycle can be summarized as follows: the sample was weighed, placed in an oven at 110 �C for 24 h, removed from the oven and placed in a desiccator, weighed again, and placed in an oven at 110 �C for 3 h. These stages were repeated until the specimen reached mass equilibrium, and only then, they were placed in the hygrothermal chamber for 60 days at 80 �C and 90 % humidity. The sample mass was periodically moni- tored, and moisture absorption (M) was calculated using M ¼ ðMw �MdÞ=Md, where Mw and Md are the wet and dry masses, respectively. In order to better understand moisture absorption and diffusion response of the composites herein analyzed, Fick’s analytical model [13], shown in Eq. (1), was applied to the data. Fig. 1 Manufacturing of a flat unidirectional laminate by filament winding 4698 J Mater Sci (2016) 51:4697–4708 123 Mt ¼ M1 1� exp �7:3 Dt h2 � �0:75 " #( ) ð1Þ where Mt is water uptake at a particular time t, M1 is the mass at a quasi-equilibrium state, D is the diffusion coef- ficient, and H is the specimen thickness. The diffusion coefficient is calculated from the absorption curve, as shown in Eq. (2): D ¼ p h 4M1 � �2 M2 �M1ffiffiffiffi t2 p � ffiffiffiffi t1 p � �2 ð2Þ where the subscripts in M and t refer to a particular mass and time, respectively. Characterization The unidirectional composites were cut longitudinally and transversely to the fiber direction. Mechanical testing was carried out using an Instron 3382 universal testing machine, with 100 and 5 kN load cells, in samples before and after the environmental conditioning. The tests were: • Tensile This test was performed at a cross-head speed of 2 mm/min in five tabbed coupons of controlled geometry dimensions following ASTM D3039-14. From this test, the elastic moduli, tensile strengths, and Poisson’s ratio were obtained. Two extensometers, one longitudinal and another transversal, were coupled to the specimens for testing and later removed just prior to rupture. • Compression The compressive strength was obtained with the combined loading compression test according to ASTM D6641-09 in five tabbed specimens. The size of the specimens was 140 9 12 9 1.4 mm3, the gage length was 12 mm, and the speed rate was 1.3 mm/min. • In-plane shear The V-notched rail shear method (ASTM D7078-12) test was chosen. The specimen (dimensions: 76 9 56 9 2.8 mm3) was centrally V-notched on both sides. The notch angle is 90� and the radius is 1.3 mm. Cross-head speed was 1.5 mm/ min and the shear modulus was determined on six samples using a strain gage rosette aligned at ± 45� at the mid-section of the sample (as seen in Fig. 2). • Short beam Interlaminar shear strength was evaluated through short beam tests, following ASTM D2344-13, using a span-to-depth thickness (s:t) ratio of 4:1. Length and width of the samples followed the recom- mendations of the standard (6 9 t and 2 9 t, respectively). • Dynamic mechanical analysis (DMA) DMA was used to evaluate the glass transition temperature (Tg) of the polymeric matrix. Analysis was carried out in a TA Instruments 2980 DMA Dynamic Mechanical Analyzer under single-cantilever bending mode at 1 Hz fre- quency, for a maximum displacement of 10 mm, under N2 atmosphere, with a heating rate of 3 �C/min, on 20 9 10 9 2.8 mm3 specimens. • Failure analysis Fractographic studies were carried out using optical micrographs in a Carl Zeiss Axio scope and scanning electron microscopy (SEM), in a Phenom ProX equipment. Failure envelope Classical failure criteria (Maximum stress, Tsai-Wu [14], Tsai-Hill [15], Hashin [16], Christensen [17], and Puck [18]) were used to predict failure envelope in the r22–s12 plane, assuming the stress in fiber direction being r1, stress in the transverse r2, and s12 the shear stress. Results and discussion Moisture uptake Figure 3 presents the mass uptake for the 0� unidirectional carbon fiber/epoxy coupons (average of five coupons) after hygrothermal conditioning. Water uptake is influenced by (i) the hydrophilic character of the matrix and fibers, (ii) fiber/matrix adhesion, (iii) micro-cracks, and (iv) voids. The resin network controls water uptake and, in turn, the absorp- tion of water influences the network. Formation of voids and micro-cracks is induced by water absorption, becoming more important as the saturation level increases [13]. Fig. 2 V-notched shear test with a bonded rosette J Mater Sci (2016) 51:4697–4708 4699 123 Hot/wet environment accelerates the deterioration pro- cess of an epoxy-based composite. Reproducibility of the experimental data is well predicted by the Fickian model, both experimental and analytical data being representative in terms of water content of the material. The initial stage till pseudo-equilibrium is typical of a thermally activated Fickian response, whereas, at longer times, the uptake occurs at a lower rate mainly due to a combination of mechanisms such as relaxation of the glassy epoxy net- work, filling of micro-voids, and debonded zones with water by wicking [19]. Thus, moisture absorption is expected to take place via diffusion. Kinetics of the diffusion process depends on tempera- ture and relative moisture absorption. The higher the rel- ative moisture absorption, the greater the absorption rate. Water saturation (&0.37 %) reached at around 30 days, and maximum mass uptake, about 0.40 %, was noticed after 42 days of conditioning, both are typical of epoxy resins. Moreover, mass uptake values were relatively low compared to literature data for the same matrix and fiber [8, 12, 13, 20], which can be justified by the high fiber volu- metric fraction and good compaction of the layers. Viscoelastic properties Figure 4 shows the storage modulus (E0) for the uncondi- tioned and conditioned composites studied. As expected, storage modulus for 0� specimens was higher than for 90� in all states (glassy, elastomeric, and rubbery). This par- ticular DMA analysis was performed under bending mode, which is markedly dependent on fiber orientation, which is stiffer for fibers aligned along the length of the specimen. Unconditioned specimens showed higher storage modulus compared to the weathered ones, which is related to the plasticizing effect promoted by moisture uptake. The aged composites showed similar behavior at the glassy state, perhaps due to a post-curing that may take place during aging. The relatively high temperatures involved, although not severe enough to break chemical bonds of the polymer [8], may contribute to the appearance of free radicals in the epoxy molecules, leading to further cross-linking [9]. The loss modulus (E00) is related to energy dissipation in the material, and composites with poor interfacial bonding are prone to dissipate more energy. Some specimens pre- sented a more flattened loss peak, where the hygrother- mally conditioned specimens had lower loss modulus, as seen in Fig. 5. Analysis of the curves shows that the behavior was more affected by fiber orientation than by weathering, and the specimens with fibers at 0o presented much higher peaks. For the 90o specimens, the aged lam- inate dissipated slightly less energy. This can be attributed Fig. 3 Mass gain obtained from experimental measurements and Fickian diffusion prediction Fig. 4 Storage modulus of the unconditioned and weathered specimens Fig. 5 Loss modulus of the unconditioned and weathered composites 4700 J Mater Sci (2016) 51:4697–4708 123 to the inhibition of relaxation processes in the composites, decreasing mobility at the fiber/matrix interface. Tan d curves are shown in Fig. 6. The unconditioned specimens showed similar temperature at the peak (i.e., Tg), and the conditioned samples showed similar and lower Tg values. The weathering effect on the composites is more clearly observed in this figure, i.e., the conditioned speci- mens show larger areas under the peak and lower Tg, suggesting matrix degradation [21]. Table 1 presents the Tg extracted from (i) intersection between the extrapolation of the elastomeric plateau and the glass state from the storage modulus curve, (ii) the loss modulus peak, and (iii) the tan d peak. The unconditioned 0o specimen showed the highest Tg, but the aged 0o spec- imen showed the lowest value in all methods; thus this specimen is more damaged by aging, confirming that fiber orientation has little effect on the glass transition of the polymer. Mechanical properties Figure 7 presents the typical load versus displacement curves for all samples. Based on the mean results of five specimens with acceptable failure, conditioning reduced tensile strength of the laminates from 1409 ± 131 to 1091 ± 114 MPa, around 29 %. Although the 0�-oriented specimens presented a slight flattening near the failure, these specimens presented a brittle behavior, typical for this type of loading/sample. Figure 7 also shows tensile behavior of the 90o-oriented composites. Brittle behavior was dominant for both non- aged_[90]4 and aged_[90]4 specimens and the ultimate load was significantly low compared to longitudinally oriented laminates. In addition, aging reduced in about 25 % their tensile strength. Even though the load versus displacement curves somewhat differ for the various longitudinally oriented specimens, all coupons showed similar failure mode (see Fig. 8a), with a sudden failure of the fibers. All 10 fractures (five for aged and five non-aged) were characterized as explosive gage middle. The failure mode for all 90o sam- ples, conditioned and non-conditioned, was also similar— lateral gage middle failure (see Fig. 8a), with only one lateral at the top tab failure. Since the optical micrographs indicated a good quality laminate, essentially free of voids (Fig. 8c, d), a weakened fiber/matrix interface may have appeared after aging (Fig. 8d), corroborated by some resin- rich areas due to water uptake and more brittle fracture at the carbon/epoxy interface, which reduced tensile strength. Furthermore, Fig. 8d suggests that broken fibers tend to initiate failure in adjacent fibers. Stress versus strain data for the tensile testing of non- aged and aged composites are shown in Fig. 9. The final properties may be summarized as follows: (i) Non-aged: E1,t = 129.8 ± 5.6 GPa and E2,t = 9.1 ± 0.5 GPa, and Fig. 6 Tan delta of the unconditioned and weathered specimens Table 1 Tg determined by different methods Specimen Tg (�C) from onset E0 drop Tg (�C) from E0 0 peak Tg (�C) from tan d peak Non-aged_[0]4 78.3 91.0 98.1 Aged_[0]4 35.2 47.7 57.9 Non-aged_[90]4 68.9 84.8 94.7 Aged_[90]4 44.7 57.8 67.4 Fig. 7 Typical load 9 displacement tensile curves for the studied families of unidirectional laminates J Mater Sci (2016) 51:4697–4708 4701 123 (ii) Aged: E1,t = 119.7 ± 4.1 GPa and E2,t = 6.3 ± 0.8 GPa. A more significant drop in E2 may be justified con- sidering that this property is more influenced by the char- acteristics of the matrix, which is more severely affected under hygrothermal aging. Figure 10 presents the longitudinal (e1) and transverse (e2) ultimate strains of the composites. For the 0� laminates, a more linear shape of the curves is noted, since e1 is much higher than e2, leading to a more consistent relationship between strains. The measured major Pois- son’s ratio (m12) was 0.31 ± 0.02 for the non-aged_[0]4 coupon. For unidirectional laminates, m21 is not easy to measure because of the very low strain magnitude. Nev- ertheless, the experimentally measured m21 was 0.020 ± 0.001, very close to that estimated through the correlation m21 ¼ m12 � E2=E1, i.e., 0.021. Compressive results for all samples are presented in Fig. 11. A non-linear behavior was noticed for non-aged and aged specimens, mainly attributed to micro-buckling caused by delaminations, as suggested by the tiny discon- tinuities along the testing direction, mainly for non- aged_[0]4 and non-aged_[90]4 coupons. It is important to bear in mind that the compressive load is transmitted though shear from the fixture to the specimen, and shear has a significant influence up to around 60 % of the test, when a purer compression stress state is achieved. The specimen with longitudinally aligned fibers is more affec- ted by aging since the decrease in matrix modulus favors micro-buckling. On the other hand, the 90o specimens are less affected by aging since, under compressive loading, failure occurs at the transversally oriented carbon fiber. Fig. 8 Optical micrographs are shown for the fractured non-aged_[0]4 a, aged_[0]4 b, non-aged_[90]4 c, aged_[90]4 d tensile coupons Fig. 9 Stress 9 strain profile plots for all coupons measured by longitudinal and transversal analogical extensometers 4702 J Mater Sci (2016) 51:4697–4708 123 Unlike tensile failure, the compression fracture surfaces are not usually normal to the loading direction, but slightly angled, since compression leads to a small degree of in- plane shear during fracture [22]. Since the laminate with compression fracture is still able to carry load after ulti- mate failure, minor damages such as longitudinal splits and delaminations may occur [23]. These characteristics can be confirmed in Fig. 12, highlighted as an implication of the fractured non-aged_[0]4 specimen. The morphology of the failure points that band of fibers fail at similar lengths, leading to the formation of kinking bands. Figure 13 displays the shear stress 9 displacement curves obtained in the V-notched rail test method. These plots show typical shear behavior, and after around 50 % of the test, most specimens started to delaminate, followed by horizontal Fig. 10 Longitudinal 9 transversal strains for determining the major (a) and minor (b) Poisson’s ratio for non-aged_[0]4 a and non-aged_[90]4 b tensile specimens Fig. 11 Typical load 9 displacement compressive curves for all specimens studied Fig. 12 Optical micrographs for the fractured non-aged_[0]4 a, aged_[0]4 b, non-aged_[90]4 c, aged_[90]4 d compressive coupons J Mater Sci (2016) 51:4697–4708 4703 123 micro-cracks at the gage area. Good compaction of the lam- inates manufactured by FW can be inferred from this graph. The breaking load was c.a. 40 % lower for conditioned specimens, indicating a strong aging effect since this is also a matrix-dominated property. The water absorbed by epoxy-based composites causes reversible plasticization of the matrix, and the combined moisture and temperature action yields dimensional changes and induces stresses that are detrimental to the fiber–matrix interface. The shear stress 9 strain profile curves are presented in Fig. 14. For both non-conditioned (Fig. 14a) and conditioned (Fig. 14b) specimens, the curve profileswere very linear,with R2[ 0.99 for all specimens. Shear modulus reduced after aging, from G12 = 5.44 ± 0.48 GPa to G12 = 3.89 ± 0.19 GPa. The failure mode of all specimens (Fig. 15) was very similar, being characterized as horizontal cracking between the notches (HGN) and multi-mode cracking at the gage section between the notches (MGN). Crack started hori- zontally close to the top notch, probably due to some stress concentration caused by machining of the notches. Later, many delamination cracks perpendicular to the loading direction appeared at the mid-section of the specimen. Representative micrographs of the notch areas of 0o spec- imens are highlighted in Fig. 15. Matrix cracking and fiber/matrix debonding dominated the failure mode for unconditioned specimens, whereas large delaminations initiated from hygrothermally damaged resin-rich regions dominated for conditioned specimens. These cracks origi- nated primarily from fiber splitting at a particular lamina and interlaminar shear at the carbon/epoxy interface. In parallel, as aging deteriorates the epoxy matrix, it decreases the overall interfacial strength of the laminate. Short beam strength (SBS) results are shown in Fig. 16. SBS values for 0o samples were 69.9 ± 1.1 and 49.5 ± 1.3 MPa for non-aged and aged, respectively. For transversally aligned fibers, SBS reduced from 12.6 ± 0.6 to 7.4 ± 0.7 MPa for non-aged and aged laminates, respectively. Short beam is one of the most recommended tests to investigate fiber/matrix interface and, from these results, it is clear that aging damages the interface, reducing interfacial strength in about 41 and 70 % for non- aged_[0]4 and non-aged_[90]4, respectively. Typical load versus displacement curves suggest that the 0� coupons failed by shear, presenting a non-linear behavior followed by discontinuities along the test, typical of delaminations and interlaminar fractures. On the other hand, the 90� samples presented a more linear behavior with brittle fracture, suggesting a strong influence of tensile load in the test. As in Almeida Jr. et al. [24], the SBS is strongly dependent on fiber orientation and it decreases from 0� to 90�. Fig. 13 Typical shear stress 9 displacement curve profiles for non- aged a and aged b coupons Fig. 14 Shear stress 9 strain curves for non-aged a and aged b specimens for the determination of shear modulus 4704 J Mater Sci (2016) 51:4697–4708 123 Even though shear is the dominant loading in this test, the internal stress state is complex and a variety of failure modes can occur, such as interlaminar shear, micro- buckling, flexure, tension, compression, or their combi- nation. In this test, the maximum shear stress occurs in a region where other stresses might also exist [25] and stress distribution may be distorted, with peak stresses taking place close to the loading points, even if the recommended span-to-thickness ratio of 4:1 [26] was used. Also, the local stress state due to the applied load may involve several shear-stress concentrations combined with transverse and in-plane compressive stresses. Figure 17 presents representative SEM images for the unconditioned 0� and 90� coupons. Failure in both 0� specimens was as expected, that is, large delamination at the mid-plane of the specimen thickness and multiple translaminar horizontal cracks, characterizing interlaminar shear failure [25, 27]. On the other hand, both 90� speci- mens showed a matrix-dominated failure and did not pre- sent evidence of shear failure. Indeed, a vertical crack initiated at the specimen bottom surface between the sup- porting cylinders, characterizing failure by flexure and primarily transverse tensile, justified by the brittle behavior of epoxy and fiber orientation. Compressive failure was not noticed. The achieved material properties are compiled in Table 2. The data presented low scatter in general, which is attributed to the advanced and automated manufac- turing process used and the laminate consolidation by hot compression prior to curing. The final laminates were nearly free of voids, as confirmed by optical and SEM micrographs. Reduction in all properties was noticed for the aged laminates due to the aforementioned reasons. Failure envelopes Failure envelopes for both non-aged and aged laminates were produced, based on several failure criteria as shown in Fig. 18, in r22–s12 plane. In most cases, matrix failure implies catastrophic failure of the laminate. In those cases, ultimate failure stress and strain of the laminates are the stress/strain corresponding to matrix failure [27]. For the maximum stress failure criteria, there is no shear. For all other criteria, shear stresses are taken into account, being more theoretically accurate for a case in which the r22–s12 plane is under study. Under tensile stress, all failure criteria tend to produce similar values due to the low magnitude of the transverse tensile strength of the unidirectional laminate. However, for the combination of transverse compression and in-plane shear, the envelopes present different predictions. In particular, the Puck criterion can be related to the previ- ous optical and SEM results, since it suggests angular fracture for the specimens under compressive loading due to a combination of micro-buckling, compression, and shear forces, typical of the Mode B inter-fiber failure. In addition, for the in-plane shear tests, delaminations, fiber/matrix debonding, and interlaminar fractures were reported, being characteristic of the Mode C inter-fiber failure in Puck criterion. Fig. 15 Representative optical micrographs from the non-aged a and aged b [0]4 tested shear specimens Fig. 16 Short beam stress versus displacement for all family of laminates studied J Mater Sci (2016) 51:4697–4708 4705 123 Conclusions In this work, the effect of weathering on tensile, com- pressive, in-plane shear, interlaminar, and viscoelastic properties of flat carbon fiber/epoxy composites manufac- tured by FW were evaluated. The composite coupons were subjected to an accelerated hygrothermal conditioning for 60 days under 80 �C and relative humidity of 90 %. Moisture absorption achieved a maximum of 0.42 % in weight, with saturation at around 33 days. Non-Fickian kinetics was found to govern moisture absorption in the final stages of weathering probably due to the leaching of low-molecular weight molecules. Matrix plasticization and a strong variation in Tg with aging were also noticed, the lowest Tg being found for the conditioned 90o specimen. Regarding mechanical properties, both tensile and compressive strength/modulus values were found to decrease for aged laminates, as well as the Poisson’s ratio. The 90o specimens were more strongly affected by aging, since a matrix-dominated failure is expected in this direc- tion. Mean shear modulus and strength reduced in as much as 30 and 40 %, respectively, and the dominant failure mode of the V-notched specimens was primarily fiber/matrix debonding and delaminations. Short beam tests Fig. 17 SEM micrographs of the fractured short beam specimens: non_aged_[0]4 (a) with a zoom at the delamination (b), and non_aged_[90]4 (c) with a zoom at the crack (d) Table 2 Summary of the measured properties Elastic Non-aged Aged Variation Strength Non-aged Aged Variation (%) E1,t = E1,c (GPa) 129.3 ± 3.6 119.7 ± 4.1 ;7 % Xt (MPa) 1409.9 ± 131.6 1091.0 ± 114.0 ;23 E2,t = E2,c (GPa) 9.11 ± 0.49 6.3 ± 0.8 ;31 % Yt (MPa) 42.5 ± 3.2 31.8 ± 1.2 ;25 G12 (GPa) 5.44 ± 0.48 3.89 ± 0.19 ;28 % Xc (MPa) 764.1 ± 79.6 621.2 ± 81.0 ;19 Yc (MPa) 134.5 ± 7.8 121.6 ± 9.8 ;10 m12 0.322 ± 0.023 0.331 ± 0.011 :3 % S12 (MPa) 68.9 ± 5.1 41.4 ± 4.0 ;40 SBS [0]4 (MPa) 69.9 ± 1.1 49.5 ± 1.3 ;29 m21 0.022* 0.029* – SBS [90]4 (MPa) 12.6 ± 0.6 7.4 ± 0.7 ;41 * Calculated 4706 J Mater Sci (2016) 51:4697–4708 123 were considered successful since the specimens failed by interlaminar shear, with large delaminations and multiple horizontal cracks throughout the specimen half-thickness for the 0o laminates. 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