The deformation and failure micromechanisms of a hybrid 3D woven com- posite were studied in tension. Plain and open-hole composite coupons were tested in tension until failure in the ll and warp directions, as well as ber tows extracted from the dry fabric and impregnated with the matrix. The macroscopic evolution of damage in the composite coupons was assessed by means of periodic unloading-reloading (to obtain the elastic modulus and the residual strain), whereas the microscopic mechanism were established by means of X-ray computed microtomography. To this end, specimens were pe- riodically removed from the mechanical testing machine and in ltrated with ZnI-containing liquid to assess the main damage modes as a function of the applied strain. The experimental observations and the predictions of an isos- train model were used to understand the key factors controlling the elastic modulus, strength and notch sensitivity of hybrid 3D woven composites in tension. Keywords: A. Polymer-matrix composites, A. Hybrid, A. 3-Dimensional reinforcement, X-ray microtomography 1. Introduction Fiber-reinforced polymers stand out among the most successful struc- tural materials due to the combination of low density, high sti ness and Email address: javier.llorca@imdea.org (J. LLorca1;2) Preprint submitted to Composites A August 5, 2013 strength and reasonable damage tolerance and impact resistance. In addi- tion, their performance can be optimized for speci c applications by choosing the ber type, volume fraction and orientation, stacking sequence, etc., nor- mally within the framework of 2D laminates. However, more sophisticated strategies are available to further improve the composite behavior for speci c applications. They include hybrid composites containing two or more ber types and 3D ber architectures. Hybrid composites in which carbon bers are used in combination with higher strain-to-failure bers have consistently demonstrated better dam- age tolerance under impact, reduced notch-sensitivity and improved fracture toughness than their carbon- ber counterparts [1, 2, 3, 4, 5, 6]. Fiber ar- chitecture has an even stronger in uence on the mechanical response [7]. The baseline materials for comparison are conventional multidirectional lam- inates manufactured from unidirectionally-reinforced pre-preg sheets. They show excellent in-plane properties due to the lack of crimping, but low delam- ination resistance under out-of-plane loads. This problem can be partially overcome with 2D fabrics, although the in-plane sti ness and strength are reduced, while non-crimp fabrics provide properties in between these ber architectures. More radical di erences (and a much wider design space) is achieved when textile manufacturing techniques are applied to create 3D ber preforms. The huge variety and the complexity of the 3D ber architecture makes the prediction of mechanical properties of 3D composites a challenging task, as their deformation and failure mechanisms are very complex and can show large di erences as a function of the loading conditions and of the 3D ber preform. Early studies [8, 9] showed that 3D woven composites presented higher failure strains than conventional multiaxial laminates, together with lower notch sensitivity and higher work of fracture. Following this pioneer work, there were many papers available in the literature on the mechanical response of 3D composites, but detailed studies focussing on the comprehen- sive assessment of the dominant damage micromechanisms in di erent 3D ber preforms have appeared recently. For instance, Gerlach et al. [10] stud- ied the e ect of the volume fraction of through-the-thickness binder on the in-plane and out-of-plane properties of 3D woven composites as a function of the strain rate. Bogdanovich et al. [11] and Ivanov et al. [12] carried out a comprehensive experimental analysis of the elastic constants and damage micromechanisms of 3D non-crimp orthogonal woven composites loaded in 2 tension. They showed that non-crimp 3D orthogonal woven composites have signi cantly higher in-plane strengths, failure strains and damage initiation thresholds than the 2D woven laminated counterparts [12]. The e ect of 3D reinforcement on delamination was studied by Pankow et al. [13], whereas Seltzer et al. [14] performed a detailed investigation on the damage mecha- nisms under low-velocity impact of various 3D orthogonal woven composites and highlighted the di erences between those observed in 2D woven materi- als. This latter study demonstrated that state-of-the art 3D characterization techniques, such as X-ray microtomography (XCT), are extremely useful to understand the initiation and progression of damage in these materials in which failure processes are inherently 3D. While the analysis of the failure mechanisms in hybrid or 3D composites has been studied in the past, there are no comprehensive studies on this topic for hybrid 3D composites containing two or more ber types. This was the main goal of this investigation, which presents a detailed analysis of the failure micromechanisms in tension of a hybrid 3D orthogonal woven composite. The evolution of damage as a function of the applied strain was monitored by means of progressive reduction in sti ness and the increment of the residual strain as well as by XCT. The composite properties (sti ness, strength) were compared with predictions obtained from the actual properties of the ber tows, which were measured independently. All this information provides a comprehensive picture of the e ect of ber hybridization, 3D ber architecture, crimping and damage on the mechanical behavior in tension of hybrid 3D woven composites. 2. Material A at composite panel was manufactured by vacuum infusion of an epoxy- vinylester resin (Derakane 8084) into a hybrid 3D orthogonal woven compos- ite. Both the dry perform and composite panel were provided by 3TEX, Inc. (Cary, North Carolina, USA) with the commercial name p3w-d00001-hx21. The preform was non-symmetric and consisted of three warp (0) and four ll (90) ber layers stacked as a cross-ply laminate [90 c,0c,90c=s2,0s2,90s2,0s2,90s2]. The schematic of the 3D ber preform is shown in Fig. 1. The bers in each layer were distributed in yarns rectangular in shape. The top four layers were made up of S2 glass bers and the bottom 2 layers of AS4C carbon bers. The hybrid layer (containing glass and carbon bers) oriented in the ll di- rection was located between the glass and the carbon layers. Each tow of this 3 Figure 1: Schematic of the unit cell of the hybrid 3D woven ber preform. Carbon ber bundles are shown in grey (dark grey for the warp direction and light grey for the ll), hybrid bundles in white and glass ber bundles in red (dark red in the warp direction and light red in the ll direction). PE z-yarn binders in the warp direction are plotted in navy blue. layer contained both AS4C and S2 glass bers, which were not intermingled but separated in two di erent zones of the tow (i.e. one half of the tow was formed by carbon bers and the other half by glass bers). It should also be noticed that every other tow was missing in the carbon layer oriented in the warp direction. In addition, the composite panel was reinforced in the through-thickness direction by z-yarn binders made up of ultra-high molec- ular weight polyethylene (PE) ber (Dyneema SK75) that went from top to bottom layers in the warp direction. Note that consecutive z-yarns were in antiphase . The nominal thickness of the dry fabric was 3.02 mm and its areal density was 4.24 kg/m2. The nominal thickness of the composite was 4.1 mm, with an areal density of 6.44 kg/m2. The overall ber volume fraction in the composite was 47% and the porosity, as measured by XCT, was 11.6%. The volume fraction of each type of ber in each direction (warp or ll) is shown 4 Table 1: Density ( ) and volume fraction of matrix and bers as a function of ber type and orientation within the hybrid composite. Material  warp ll total (g/cm3)(%) (%) (%) Glass S2 2.48 15.4 14.7 30.1 Carbon AS4C 1.78 4.3 10.7 15.0 Polyethylene SK75 0.97 1.9 | 1.9 Total bers 21.6 25.4 47 Matrix 1.02 53 in Table 1. It was determined from the areal density of the individual plies and the ber density, which were provided by the manufacturer. In addition, the matrix volume fraction was determined from the matrix density and the weight of the composite panel before and after in ltration and this is also given in Table 1. 3. Experimental Techniques 3.1 Mechanical characterization Rectangular specimens of 250 254.1 mm3were machined from the plate with the longest dimension aligned in either the warp or ll direction for the mechanical tests. Cu tabs of 50 mm in length were glued to the specimens, leading to a free length of 150 mm. They were tested in tension in an electromechanical universal testing machine (Instron 3384) following the recommendations of the ASTM Standard D3039 [15] and of the ASTM Standard D5766 [16] for rectangular coupons with a central hole of 4.1 mm and 11 mm in diameter. Tests were carried out under stroke control at 2 mm/min and the load was continuously measured during the test with a load cell of 150 kN. Since composite is non-symmetric (and coupling between bending and extension might occur), the longitudinal strain was recorded on both faces of the specimen using an extensometer of 50 mm gage length on one face and digital image correlation (Vic2D) on the other. Periodic unloading-reloading was carried out in one test in each direction (warp or ll) to estimate the sti ness degradation as a function the applied strain. In addition to the composite specimens, the mechanical properties of the di erent ber tows were also measured following the recommendations of 5 the ASTM Standard D4018 [17]. To this end, individual tows of carbon, glass and PE were carefully extracted from the dry fabric to avoid damage. They were impregnated with Derakane 804 epoxy-vinylester resin and cured at room temperature, following the instructions of the resin manufacturer, with a geli cation time of 30 min. The linear density of the ber tows before and after impregnation was measured. Glass ber tabs were glued to the impregnated tows, leading to a free length of 150 mm. They were tested in tension until failure under stroke control at 5 mm/min. Load and strain were continuously recorded during the test, the latter with an extensometer. Thermo-gravimetric analyses of the impregnated carbon and glass yarns were conducted in a TGA Q50 Thermal Analyser (TA Instruments). All ex- periments were carried out in air with a purge rate of 60 ml/min and a heating rate of 10C/min up to 1000C. The sample weight was approximately 10 mg and 18 mg for carbon and glass ber tows, respectively. The ber weight fractions were 64% and 69% for carbon and glass bers, respectively, which led to ber volume fractions of 50.5% and 48%. This methodology could not be used to obtain the ber volume fraction in the PE ber yarns because of the low melting point of the bers. 3.2 Damage inspection The damage micromechanisms in the composite specimens tested in ten- sion in the warp and ll directions were analyzed as a function of the applied strain by means of XCT. To this end, tow coupons in each direction were deformed up to pre-de ned strains. The tests were stopped and the speci- mens were immersed in a liquid for 2 hours while holding the displacement constant to facilitate the penetration of the liquid into the cracks. The liq- uid was composed of 60 g of ZnI in 10 ml of water, 10 ml of ethanol and 10 ml of Kodak Photo-Flo 200. The specimens were removed from the ma- chine and inspected by XCT as detailed below. Previous investigations have demonstrated that the immersion in this liquid does not modify the damage mechanisms in composites [18]. XCT of the specimens was performed with a Nanotom 160NF (Phoenix) at 70-90 kV and 120-180 mA using a tungsten target. 1800 radiographs were acquired for each tomogram with an exposure time of 1000 ms. The 3D to- mograms were reconstructed from the radiographs using an algorithm based on the ltered back-projection procedure for Feldkamp cone beam geome- try. The reconstructed volumes presented a resolution of 13-15 m/voxel 6 0500100015002000250030003500 00.010.020.030.040.05AS4C CarbonS2 GlassStress (MPa)Strain(a) 0500100015002000250030003500 00.010.020.030.040.05SK75 PolyethyleneStress (MPa)Strain(b)Figure 2: Stress-strain curves of the ber tows. (a) Carbon and glass ber tows. (b) Polyethylene ber yarns. and were analyzed using the commercial software VGStudio Max. Emphasis was placed on the qualitative assessment of the main damage mechanism in each material as a function of applied strain rather than in the quanti ca- tion of speci c damage modes (delaminated surface, volume of cracked tows, fraction of broken bers, etc.). 4. Mechanical properties 4.1 Fiber tows Six tows were tested in tension until failure for each ber type and the average stress in the bers is plotted as a function of the applied strain in Fig. 2 for the carbon (AS4C), glass (S2) and polyethylene (SK75) bers. The matrix contribution was neglected in the stress analysis and the average stress in the carbon and glass bers was computed from the applied load P, the tow cross-section and the ber volume fraction, as determined by the thermo-gravimetric analysis. In the case of the PE yarns, the average ber stress was computed as Pf=y, wherefstands for the ber density and y is the linear density of the dry yarn. The carbon tows (Fig. 2a) presented a linear elastic behavior up to 0.5% strain. Beyond this point, three specimens showed a slight reduction in sti - ness whereas two maintained the initial modulus until failure. One specimen 7 Table 2: Elastic modulus, tensile strength and fracture strain of the bers within the ber yarns Fiber type Elastic modulus Tensile strength Failure strain (GPa) (GPa) (%) Glass S2 821 2.40.3 4.01.0 Carbon AS4C 2072 2.80.4 1.50.2 Polyethylene SK75 851 2.50.2 4.10.2 presented evidence of damage (sudden stress drop) before nal failure. All of them failed in a brittle fashion when the strain was in the range 1.2% - 1.9 %. Evidence of carbon sti ening was not found in these tests. The glass ber tows (Fig. 2b) also presented a linear elastic behavior up to 1.6%. The elastic modulus decreased with the applied strain beyond this point and damage prior to failure was observed in most of the yarns for applied strains in the range 2% to 3%. Final fracture occurred in a brittle fashion, as in the case of the carbon ber tows, although at higher strains (in the range 3% to 5%). The PE ber yarns (Fig. 2c) showed a more ductile behavior and the stress-strain curves were non-linear in the whole range of deformation. The elastic modulus increased up to 1.0-1.5% strain and then decreased up to fail- ure, which occurred at an applied strain in the range 4.0%-4.3%. It should be noted that the scatter in the stress-strain curves of the PE yarns was very small. The average elastic modulus, failure strain and tensile strength of the ber tows, together with the corresponding standard deviations, are summarized in Table 2. 4.2 Tensile properties of plain coupons Three tensile tests were carried out in the ll and warp directions and the corresponding stress-strain curves are plotted in Figs. 3a and 3b, re- spectively. The strains were measured by conventional extensometry on the glass face of the coupon and by digital image correlation on the carbon face. They are plotted in Fig. 3c for two representative tests in the ll and warp direction. The di erences in the strain between both faces were negligible, indicating that the extension-bending coupling induced by the asymmetry of the composite did not play a signi cant role in the tensile behavior at least until the fracture of the carbon bers. The stress-strain curves in the ll direction (Fig. 3a) presented a linear- 8 0100200300400500600 00.010.020.030.04Stress (MPa)Strain(a)fill directionfill direction 0100200300400500600 00.010.020.030.04Stress (MPa)Strain(b)warp direction 00.0050.010.0150.020.0250.03 050100150200250300350warp (C side, DIC)warp (glass side)fill (C side, DIC)fill (glass side)Strain Time (s)(c) C fiber failureFigure 3: Stress-strain curves of the plain composite coupons. (a) Fill direction. (b) Warp direction. The solid circles in the curves indicate the instants in which the test was stopped and the coupons were examined by means of XCT. (c) Tensile strains measured by conventional extensometry on the glass side and by digital image correlation on the carbon side for two representative tests in the ll and warp direction, respectively. 9 elastic behavior up to a strain in the range 1.2% to 1.6%, which is similar to the failure strain of the carbon ber tows (Fig. 2a). There was a sudden drop in the load bearing capacity of the composite but the load increased slightly with further strain until catastrophic failure, which occurred at an applied strain below 2.5%. The initial part of the stress-strain curve was also linear in the warp direction (Fig. 3b), but a slight non-linearity was detected before the load peak, which occurred when the applied strain was in the range 1.3% to 1.7%, as in the ll direction. Nevertheless, the load bearing capacity of the composite increased with further straining and the maximum strength was attained for applied strains in the range 2.1% to 3.1%, just before the coupon failed catastrophically. This failure strain is compatible with the development of damage in the glass ber tows. Thus, composite coupons loaded in tension along the ll and warp di- rection presented two peaks. The maximum strength along the ll direction was attained in the rst peak and it was controlled by the fracture of the AS4C carbon ber yarns. The maximum strength in the warp direction was attained in the second peak and it was controlled by the fracture of the S2 glass ber tows. It should be noted that the main di erences in the compos- ite architecture between the ll and warp direction were the presence of an extra hybrid ply with the bers oriented in the ll direction and that the PE z-yarn binders were oriented along the warp direction. The evolution of the sti ness as a function of the applied strain was obtained from the tests in which the coupon was periodically unloaded and reloaded. It is plotted in Figs. 4a and 4b for the coupons deformed in the ll and warp directions, respectively. The sti ness was obtained from the slope of the stress-strain curve at the beginning of the unloading cycle, from the initial unloading point up to the point at which the load has decreased by 20%. The residual strain (understood as the longitudinal strain at zero load at the end of the unloading part of the cycle) is also plotted in these gures. In the case of the coupons deformed in the ll direction (Fig. 4a), the sti ness remained constant (or even increased slightly) up to the onset of carbon ber failure (1.3%-1.4%), even though the continuous increment of the residual strain in this range indicated the development of irreversible damage mechanisms in this region. It should be noted that the PE z-yarn oriented in the warp direction induced ber crimping in the ll tows and the straightening of the carbon and glass ber tows upon deformation could compensate the reduction in sti ness due to damage, leading to this plateau 10 010203040 00.050.10.150.20.250.3 00.511.522.53StiffnessResidual strainLongitudinal stiffness (GPa)Residual strain (%) Strain (%)(b) warp direction010203040 00.050.10.150.20.250.3 00.511.52Stiffness ResidualstrainLongitudinal stiffness (GPa)Residual strain (%) Strain (%)(a) fill directionFigure 4: Evolution of the elastic modulus and of the residual strain (at zero load) with the applied strain. (a) Fill direction. (b) Warp direction. in the longitudinal sti ness. Fracture of the carbon ber yarns led to a dramatic reduction in sti ness prior to failure while the permanent strain increased accordingly. In the coupons oriented along the warp direction (Fig. 4b), sti ening of the coupon due to the straightening of glass and carbon ber yarns was not signi cant and a continuous reduction in sti ness was observed from the beginning of the test, together with an increase in the residual strain. Both (reduction in sti ness and increase in residual strain) were enhanced after fracture of the carbon ber tows. 4.3 Tensile properties of open-hole coupons The tensile stress-strain curves of the open-hole tests are plotted in Figs. 5a and 5b for coupons with a central hole of 4.1 mm and 11 mm in diame- ter, respectively. These curves present similarities and di erences with those measured in plain coupons (Fig. 3). The maximum strength was also at- tained at lower strains when the composite was loaded in the ll direction and was dictated by the fracture of the carbon bers. The coupons loaded in the warp direction attained the maximum strength at higher strains after the failure of carbon bers. Nevertheless, damage in the ligament induced by the presence of the hole leads to a non-linear stress-strain curve even in the samples loaded in the ll direction. The non-linearity increased with the hole diameter. Moreover, the specimen failure after the peak strength was gradual, particularly in the specimens with large holes (Fig. 5b). This 11 behavior is indicative of a ductile failure mechanism in which the material in the ligament failed progressively as the applied strain increased. The notch sensitivity of the hybrid 3D woven composite in the ll and warp directions is shown in Fig 5c, in which the failure stress (normalized by the average failure stress of the plain coupon) is plotted as a function of the hole diameter (divided by the coupon width). The experimental results for the open hole tests with di erent hole diameter are very close to the straight line which indicates that failure strength was proportional to the ligament width (particularly in the ll direction) and thus that the stress concentration induced by the circular hole did not in uence the failure strength. Reduced notch sensitivity in carbon/epoxy 3D woven composites was already reported [8, 9] and it was partly attributed to the presence of geometrical aws that were broadly distributed in strength and space; and partly to the coarseness of the reinforcing yarns, which led to extensive debonding and reduced stress intensi cation around sites of failure. Nevertheless, the hybrid 3D woven composites presented better notch insensitivity, indicating that hybridization with glass and PE bers improves the damage tolerance. 5. Damage and failure micromechanisms 5.1 Loading along the ll direction XCT provides unique information regarding the microscopic damage and failure micromechanisms which develop during deformation of the complex 3D microstructure of the hybrid composite material. XCT was carried out in one coupon deformed in the ll direction and deformed up to strains of 1.05% and 1.74% (marked with solid circles in Fig. 3a). The rst analysis occurred in the linear region of the stress-strain curve, before carbon ber breakage took place. Two cross-sections of the investigated volume parallel to the loading axis are shown in Figs. 6a and 6b. Bright regions in the tomograms correspond to cracks (or pores), whose contrast was enhanced by the presence of ZnI. The two sections in Fig. 6 are parallel and one of them (Fig. 6b) contains the PE z-yarns, which run in the vertical direction. The two carbon layers running in the ll and warp directions are at the top in Fig. 6, and the missing yarn every other yarn in the warp direction is clearly visible in Fig. 6a. The four glass ber layers are at the bottom and the hybrid layer (light grey/dark grey contrast) is in between. The resin- rich area due to the missing yarn enhanced crimping of the adjacent carbon and carbon/glass yarns in the ll directions, leading to delamination of the 12 0100200300400500 00.010.020.030.04fillwarpStress (MPa)Strain(a)4.1 mm hole diameterfill direction 0100200300400500 00.010.020.030.04fillwarpStress (MPa)Strain(b)11 mm hole diameterfill direction 00.20.40.60.81 00.20.40.60.81FillWarpNormalized strengthHole diameter / coupon width(c)fill directionFigure 5: Stress-strain curves of the open-hole composite coupons in the ll and warp direction. (a) hole diameter of 4.1 mm. (b) Hole diameter of 11 mm. (c) Notch sensitivity of the 3D hybrid woven composite. The notch insensitive behavior is given by the dashed straight line. 13 ll tows from the matrix (marked with arrows). These delamination cracks propagated along the tows and notice that delamination was asymmetric: only the upper or the lower tow boundary was delaminated from the ma- trix. In addition, tensile transverse cracks were observed in the ber tows oriented in the warp direction (perpendicular to the tensile loading). These cracks propagated through the matrix (and along the interface) within the ber tows. The parallel cross-section in Fig. 6b depicted the same failure mechanisms (namely, delamination of the carbon and hybrid ll tows from the matrix near the missing warp tows and tensile cracking of the warp tows) but it is worth noting that the PE z-yarns arrested the propagation of the de- lamination cracks. In addition, delamination of PE z-yarns was also evident (marked with arrows). More information about the damage mechanisms was obtained from the cross-sections perpendicular to the loading axis, which are shown in Fig. 7. The delamination cracks at the boundary of the carbon and glass ber tows (already shown in Fig. 6) are also visible in these tomograms, together with delamination cracks at the boundary of PE z-yarns (Fig. 7a). In addition, cracking within the ber tows parallel to the loading axis (marked with ar- rows) was found (Figs. 7a and 7b). Their number was higher in the ber tows surrounded by the PE z-yarns (Fig. 7a), in which cracks oriented in the horizontal and vertical direction are seen. The cracks parallel to the loading axis within the ber tows could not be attributed to the far- eld tensile stress eld and were due to the longitudinal shear stresses that develop within the tow as a result of asymmetric tow delamination from the matrix, i.e. the up- per boundary of the tow delaminates from the matrix and deforms along the loading direction while the lower boundary remains adhered to the matrix (orviceversa ). The second XCT inspection of the specimen loaded in the ll direction was carried out at 1.74% strain after fracture of the carbon ber tows. The tomograms in the longitudinal sections parallel to the loading axis are shown in Fig. 8. Fracture of the carbon ber tows is clearly visible in the upper re- gion of the specimen as well as extensive delamination of the ber tows from the matrix (Fig. 8a). In addition, delamination cracks propagated from one layer to another through the warp tows. The parallel section containing the PE z-yarns (Fig. 8b) shows that they were able to arrest the longitudinal delamination cracks. In addition to these failure mechanisms, tensile cracks within the warp were found in both longitudinal sections. The inspection 14 Figure 6: Tomograms of the coupon loaded in the ll direction up to 1.05% strain (a) Longitudinal section parallel to the (horizontal) loading axis (b) Another longitudinal section parallel to the (horizontal) loading axis which contains the PE z-yarns. Glass ber tows appear light grey, carbon bers, PE z-yarns and matrix appear dark grey and white regions correspond to cracks in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: dstands for delamination cracks from the ber tows and z-yarns and t indicates transverse tensile cracks within the ber tows. 15 Figure 7: Tomograms of the coupon loaded in the ll direction up to 1.05% strain. (a) Transverse section perpendicular to the loading axis which contains the PE z-yarns. (b) Another transverse section perpendicular to the loading axis. Glass ber tows appear light grey, carbon, PE bers and matrix appear dark grey and white regions correspond to cracks in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: dstands for delamination cracks from the ber tows and z-yarns and sindicates shear cracks within the ber tows. 16 of the composite sections perpendicular to the loading axis (Fig. 9) showed that fracture of the carbon tows has enhanced the development of the de- lamination cracks at the boundary of all the ber tows (carbon, glass and PE). As a result, the number of shear cracks parallel to the loading axis have also increased and multiple shear cracks were visible within many glass ber tows (Fig. 9) 5.2 Loading along the warp direction XCT was carried out in one coupon deformed in the warp direction at strains of 1.19% and 2.15% (marked with solid circles in Fig. 3b). The tomograms of cross-sections parallel and perpendicular to the loading axis deformed up to 1.19% are depicted in Figs. 10 and 11, respectively. The main damage mechanisms observed in the longitudinal sections which did not contain PE z-yarns were ber tow delamination and tensile transverse cracks within the ber tows (Fig. 10a). Delamination cracks were occasionally de ected into the ber tows, propagated across the tow and led to another delamination crack at the opposite tow surface. The sections containing PE z-yarns also showed extensive delamination of the PE yarns from the matrix and carbon and glass ber tows, while delamination cracks between matrix and carbon or glass were arrested by the z-yarns (Fig. 10b) The XCT of the sections perpendicular to the loading axis at 1.19% strain (Fig. 11) showed the expected damage mechanisms: delamination cracks between the matrix and the ber tows (PE, carbon and glass) as well as extensive shear cracks within the carbon and glass ber tows induced by the asymmetric delamination of the ber tows from the matrix. XCT in the composite coupons loaded in the warp direction after the frac- ture of the carbon yarns are depicted in Figs. 12 and 13. Fig. 12 includes the cross-sections parallel to the loading axis, in which extensive tensile trans- verse cracks in the carbon tows appear as the main damage mechanism (Fig. 12a). Delamination cracks at the ber tow interfaces are visible but pref- erentially at the PE z-yarn boundaries (Fig. 12b). They propagated into the carbon and glass ber tows, leading to extensive intratow cracking. The cross sections perpendicular to the loading axis showed extensive damage in the carbon ber tows (Fig. 13a and 13b), while shear cracking of the glass and hybrid ber tows was visible throughout the section. Maximum shear crack density within the tows was found on the sections containing the PE z-yarns (Fig. 13b) and it was triggered by the shear stresses within the tow 17 Figure 8: Tomograms of the coupon loaded in the ll direction up to 1.74% strain. (a) Longitudinal section parallel to the (horizontal) loading axis (b) Another longitudinal section parallel to the (horizontal) loading axis which contains the PE z-yarns. Glass ber yarns appear light grey, carbon ber yarns and matrix appear dark grey and white regions correspond to cracks or voids in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: fstands for the fracture of the carbon ber tows, ttensile transverse cracks within the ber tows, ddelamination of the ber tows, scslanted translaminar cracks within the ber tows created by propagation of delamination cracks through the tow and vvoids. 18 Figure 9: Tomograms of the coupon loaded in the ll direction up to 1.74% strain. (a) Transverse section perpendicular to the loading axis (b) Another transverse section per- pendicular to the loading axis which contains the PE z-yarns. Glass ber yarns appear light grey, carbon ber yarns and matrix appear dark grey and white regions correspond to cracks or voids in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: ttensile cracks within the ber tows, ddelamination of the ber tows, sindicates shear cracks within the ber tows, and vvoids. 19 Figure 10: Tomograms of the coupon loaded in the warp direction up to 1.19% strain. (a) Longitudinal section parallel to the (horizontal) loading axis (b) Another longitudinal section parallel to the (horizontal) loading axis which contains the PE z-yarns. Glass ber yarns appear light grey, carbon ber yarns and matrix appear dark grey and white regions correspond to cracks in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: ttensile transverse cracks within the ber tows, ddelamination of the ber tows, sc slanted translaminar cracks within the ber tows created by propagation of delamination cracks through the tow. 20 Figure 11: Tomograms of the coupon loaded in the warp direction up to 1.19% strain. (a) Transverse section perpendicular to the loading axis (b) Another transverse section perpendicular to the loading axis which contains the PE z-yarns. Glass ber yarns appear light grey, carbon ber yarns and matrix appear dark grey and white regions correspond to cracks in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: sshear cracks within the ber tows, ddelamination of the ber tows, and vindicated voids (which appear black). 21 promoted by the delamination of the z-yarns. 6. Prediction of the coupon properties Although XCT analyses showed the development of di erent damage mechanisms upon tensile deformation, the tensile strength of the compos- ite coupons seemed to be controlled by the fracture of the ber yarns in the loading direction. Previous investigations of the tensile behavior of 3D woven composites [8] have shown that the tensile strength can be estimated fairly well from the stress carried by the individual tows oriented in the loading direction assuming an isostrain approach. According to this hypothesis, the composite mechanical behavior in the ll and warp directions should be pre- dicted from the experimental values of the ber elastic modulus and failure strain obtained from the ber tow tests using a simple model based on the rule of mixtures. The composite elastic modulus in the ll ( Ef) and warp (Ew) directions could be expressed as Ef=Ecff c+Es2ff s2+Em(1ff cff s2) (1) Ew=Ecfw c+Es2fw s2+Epefw pe=2 +Em(1fw cfw s2fw pe=2) (2) whereEc,Es2andEpestand for the elastic moduli of the bers (Table 2), ff c,ff s2are the carbon and glass ber volume fractions in the ll direction (Table 1), and fw c,fw s2andfw pethe carbon, glass and PE ber volume fractions in the warp direction (Table 1), respectively. This model assumes that the contribution to the elastic modulus of the ber yarns perpendicular to either the ll or the warp direction is given by the matrix modulus, Em, which was taken as 2.9 GPa, according to the manufacturer. In addition, only one half of the volume fraction of the PE bers was considered to compute the elastic modulus in the warp direction because of the shape of the PE z-yarns. The model predictions are compared with the experimental data in Table 3, showing excellent agreement and supporting the use of the isostrain approach to model the elastic deformation in tension of hybrid 3D woven composites. The prediction of the tensile strength in the ll direction, Xf, was carried out using the same model under the assumption that the composite behavior is linear until fracture. Thus, Xf=Efc (3) 22 Figure 12: Tomograms of the coupon loaded in the warp direction up to 2.15% strain. (a) Longitudinal section parallel to the (horizontal) loading axis (b) Another longitudinal section parallel to the (horizontal) loading axis which contains the PE z-yarns. Glass ber yarns appear light grey, carbon ber yarns and matrix appear dark grey and white regions correspond to cracks in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: fstands for the fracture of the carbon ber tows, ttensile transverse cracks within the ber tows, ddelamination of the ber tows, scslanted translaminar cracks within the ber tows created by propagation of delamination cracks through the tow. Table 3: Experimental results and model predictions for the elastic modulus and tensile strength of the 3D hybrid woven composite in the ll and warp directions orientation Elastic Modulus Tensile strength Experiments Model Experiments Model (GPa) (GPa) (MPa) (MPa) Fill 38.41.6 38 53142 545 Warp 24.60.7 24.6 39535 612 23 Figure 13: Tomograms of the coupon loaded in the warp direction up to 2.15% strain. (a) Transverse section perpendicular to the loading axis (b) Another transverse section perpendicular to the loading axis which contains the PE z-yarns. Glass ber yarns appear light grey, carbon ber yarns and matrix appear dark grey and white regions correspond to cracks in ltrated by ZnI. Arrows indicate the di erent damage mechanisms: fstands for the fracture of the carbon ber tows, sshear cracks within the ber tows, and d delamination of the ber tows. 24 wherecis the failure strain of the carbon ber tows (Table 2). In the case of the warp direction, the tensile strength was attained at the fracture of the glass bers after the carbon ber yarns were broken. Neglecting the contribution of the carbon ber yarns to the elastic modulus, the composite tensile strength in the warp direction, Xwcan be computed as Xw= Es2fw s2+Epefw pe=2 +Em(1fw cfw s2fw pe=2) s2 (4) wheres2stands for the fracture strain of the glass ber tows. The model predictions are compared with the experimental results in Table 3. The model predictions in the ll direction are very accurate, as this is the result of the linear behavior of the composite until the fracture of the carbon ber tows (Fig. 3a). However, the model overestimated the composite strength in the warp direction because the failure strain of the composite (2.8 0.3%) was much lower than the failure strain of the glass ber bundles (4.0 1%), which controlled the composite strength in this orientation. In fact, if the failure strain of the S2 glass ber yarns, s2in equation (4) is substituted by the failure strain the coupon in the warp direction (2.8%), the isostrain model prediction for the composite strength in the warp direction drops to 428 MPa, close to the experimental results. The early fracture of the glass ber bundles was very probably promoted by the bending stresses induced in the asymmetric hybrid laminate after the failure of the carbon bers and limited to maximum loading bearing capability of the hybrid composite. 7. Concluding remarks The analysis of the deformation and failure micromechanisms of a hy- brid 3D woven composite in tension showed that the elastic deformation was controlled by the elastic modulus and volume fraction of ber yarns ori- ented in the loading direction. Due to the complex 3D architecture, damage developed at low strains in the form of debonding of the ber tows (lon- gitudinal and z-yarns) as well as of tensile and shear cracking of the ber tows. Tow debonding was arrested by the z-yarns whereas tensile transverse cracks were concentrated within ber tows perpendicular to the loading axis. Shear cracking of the ber tows was due to asymmetric debonding of the ber tows from the surrounding material and developed in longitudinal and transversal tows. Nevertheless, the in uence of these damage mechanisms in the elastic modulus was limited and, depending on the architecture, was balanced by the straightening of the ber tows upon tensile straining. Thus, 25 the elastic moduli in both ll and warp directions was accurately predicted by an isostrain model using the actual properties of the ber tows within the composite. Elastic deformation stopped abruptly when the applied strain reached  1.5% due to the brittle failure of the carbon ber tows. This rst peak in the stress-strain curve could be adequately predicted by the isostrain model and bending stresses (induced by the non-symmetric laminate) did not in uence the mechanical response. Nevertheless, fracture of the carbon yarns did not lead to the laminate fracture because the laminate was held together by the PE z-yarns while the glass ber tows continued supporting the applied load. This led to another peak in the stress-strain curve, which was controlled by the fracture of the glass ber tows. Whether or not the second peak was higher than the rst one depended on the volume fraction of carbon and glass bers oriented in the loading direction. It should be noted that the strain at which this second peak was attained (2-3%) was signi cantly lower than the failure strain of the glass ber tows ( 4%) probably due to the development of bending stresses in the damaged coupon. Thus, the full contribution of the glass bers to the composite strength was not employed but their presence increased the fracture strain and the energy dissipated during fracture. Thus, hybridization of the 3D woven composite led to a notch-insensitive behavior as demonstrated by open-hole tests. This preliminary investigation points out the complexity of the deforma- tion and damage mechanisms in hybrid 3D woven composites and further detailed analyses are required to understand the interaction among the dif- ferent processes. This information will be critical to establishing design rules to optimize the behavior of these composites under di erent loading condi- tions, taking advantage of the di erent sti ness and failure strains of the bers in the composite and of the asymmetry in the laminate. Acknowledgements This investigation was supported by the Spanish Min- istry of Economy and Competitiveness through the project MUDATCOM (MAT2012-37552) and by the Comunidad de Madrid through the program ESTRUMAT-CM (S2009/MAT-1585). The authors acknowledge 3TEX Inc. for providing the material and the areal densities. The help of Mr. J. C. Rubalcaba, Mr. J. L. Jimenez and Ms. E. Moreno is also gratefully acknowl- edged. 26