Transverse Cracks at Ply Drops in Fiberglass Laminates

Glassfiber reinforced epoxy resin laminates containing internal ply drops were fabricated using resin film infusion.The resin in the region of the ply drops was enhanced using two types of additive: aramid pulp and vapor grown carbon nanofiber (VGCNF).Laminate lay-ups of[0/0*/0]T and[+45/-45/0*/-45/+45]T (where * indicates the ply that was discontinued) were tested in both static and fatigue loading.Failures init iated at the resin pocket formed at the termination of the dropped ply, which was detected by both edge photomicrographs and a jump in longitudinal strain at the ply drop location.For the[0/0*/0]Tlaminates, addition of the aramid pulp resulted (when compared to the non-modified epoxy-only resin) in a large increase in the quasi-static stress required to init iate damage, and a slight increase in laminate strength.Addition of the carbon nanofiber resulted in essentially no change in the stress required to initiate damage, and a reduction in the ult imate strength.The trend for the[+45/-45/0*/-45/+45]T laminates showed an increase in the damage initiat ion quasi-static stress in both the aramid pulp and VGCNF laminates, appreciably no change in ultimate strength due to the aramid pulp, and a notable ult imate strength increase due to the inclusion of the VGCNF.Results from fat igue testing showed a large fatigue penalty due to the inclusion of the ply-drop, and no significant difference in fatigue life due to the inclusion of the nano-scale rein forcements.Finite element stress analysis modeling of the crack growth indicates the effects of having increased matrix modulus in the resin pocket zone, as well as increased value of laminate critical energy release rate.The analysis indicates the initiat ion of failu re v ia transverse cracks in the resin pocket.This was confirmed by photomicrographs taken at fixed load intervals.


Introduction
The rapid growth of the global econo my as spurred a corresponding increase in the need for electric power. In addition, it is desired to have this new power at reasonable cost and with minimal carbon output (to address concerns related to carbon output and global warming). In response t o th ese needs , th e wind tu rb ine generato r ind ust ry (in clu d ing g ov ern men t and acad emic p artners ) has stepped-up research to reduce the cost, increase the life, and improve the reliability of the wind generator systems. Key components of the turbines are the long, slender rotating blades. The highly-loaded blades are typically fabricated fro m fiber-reinforced co mposite materials to reduce weight and increase the blade fatigue life [1]. Glass fiber is the most co mmon ly used reinforcing material, d ue to its h igh strength and low cost when co mpared to other co mposite reinforcing fibers. As with any wing, the blade loading is higher in the root region, near the blade hub, and becomes less toward the blade tip. Therefore, the laminate thickness (s kin , s p ar, and web ) fo r th es e lo ng , h ig h ou tp ut blade designs must necessarily change from the blade hub to the tip.Therefore, this study was initiated to focus on the damage development at ply d rops.Ply drops in this study are the locations in a laminate where an internal p ly is terminated.This termination creates an internal edge, where the plies above and below the terminated ply must co me together beyond the location of the termination, as shown in Figure 1.Th is termination naturally creates a small triangular-shaped wedge at the location of the termination, which fills with resin when the laminate is cured.A major concern is that the high stresses created in this region can initiate laminate failure, as has been demonstrated both experimentally and analytically [2][3][4][5][6][7][8]. The failu re in ply drop regions characteristically contains a transverse crack in the resin pocket, which, under increased static or fatigue loading, leads to growth of delaminations fro m the ends of this transverse crack. Therefore, addressing methods to analytically model the location and secondary crack progression, experimentally detect the transverse crack initiat ion, and delay the occurrence of these cracks in the ply drop resin pocket via the use of nano-scale resin enhancements are novel, and critical to imp roving the life and reliab ility of wind turbine generators.
To address these issues, the purpose of this study was threefold: first, to analytically explore the effect of improved mat rix modulus and interlaminar fracture toughness on the propensity of the resin pocket to crack and delaminate (and identify the crack sequence and location); secondly to create a rapid and accurate method to identify when the resin cracks fo rm in laminates under test; and finally to investigate the use of nano-enhanced resins to delay the onset of cracking and delamination in the region of the resin pocket at the ply drop.

Speci men Fabricati on
Laminate panels were fabricated with two types of lay-ups:[0/0*/ 0] T and[+45/-45/0*/-45/+45] T , where the * indicates the ply that was dropped (extended only one-half of the length of the specimens).Thin laminates were selected as a research tool to emphasis the effect of the ply drop and resin nano-scale modifications. The laminates were fabricated using resin film infusion of Epon 862 epoxy resin (with W hardener), and autoclave cured with the following cycle: ramp to 121℃(250 °F), hold for 1 hour, ramp to 177℃ (350 °F), hold for 2 hours, 0.758 M Pa (110 psi) pressure; vacuum-bagged.In several of the panels, during the panel lay-up, additional resin films strips 7.6 cm wide were placed immed iately belo w and above the ply drop region, as shown in Figure 2.These films contained 10 g/m 2 o f either the aramid Kevlar® pulp (DuPont Type 953) or Applied Sciences, Inc carbon nanofiber PR-25-PS-OX(no heat treat; functionalized).Fu rther processing details and a discussion of the benefits of the nano-scale modifications has been previously detailed [9][10].The fiberglass used was Owens Corning WindStrand Un iweave Knytex XHPCD1200-127cm and XHPDB1000-127 cm.The resulting laminates had a typical fiber volu me fraction of 47-51%.  The cured panels were tabbed and cut into 2.54 cm wide strips for tensile testing, as shown in Figure 3.The specimen tint fro m the aramid pulp (top) and carbon nanofiber (bottom) is visible.Note that the static tensile specimens received 0/90 strain gages at two locations: in the central areas of thick regions of the specimens, and just at the ply drop transitions.The purpose of the gages in the thick region was to measure the far-field strains, and the gages in the ply drop region were intended to indicate the initiat ion of damage.

Speci men Testing
The quasi-static specimens were tested in accordance with ASTM D3039 [11], at a rate of 1.27 mm/ min. under ambient laboratory conditions. The loading in several representative specimens (having polished edges) was interrupted at fixed intervals. These specimens were then removed fro m the load frames, edge photomicrographs were taken, then the specimens were returned for futher loading.
The tension-tension fatigue tests were conducted with no strain gages, at a rate of 2 Hz, and R = σ min /σ max = 0.1.

Quasi-Static Tests of[0/0*/0] T
A stress versus strain plot for one of the epo xy-only (no nano additive in the ply drop reg ion) quasi-static specimens is shown in Figure 4.It is important to note that the strain gage location for the longitudinal strain values used in these figures was directly on top of the ply drop location.The ply drop region is somewhat different fro m specimen to specimen (due to fiber wash, etc).In addition, large strain gradients were expected in this region.Therefore, the strain output is useful to indicate damage init iation and gro wth, and not useful in a direct numerical co mparison between specimens.Also, note that each specimen had a thick and a thin region, due to the ply drop across the center of each test coupon.The stress was calculated as the axial force on the specimen div ided by the cross-sectional area of the thin region (hence nominal stress values are used here).
Note, in Figure 4, the slight ju mp in the strain value when the stress reaches 270 MPa.Th is corresponds to the photomicrograph fro m a similar specimen, as shown in Figure 5.Each frame in Figure 5 shows the dropped ply, the plies adjacent to the dropped ply, and the resin wedge formed.A faintly visible transverse crack can be seen in the specimen having been removed at 316 M Pa.Since the specimens were removed at fixed load values, based on the micrographs only, we know the init ial crack formed between 0 and 316 MPa.A lso note the further damage development due to a subsequent load increase, seen as mu ltiple cracks in the micrograph taken at 765 MPa stress level.The final frame in Figure 5 shows the laminate edge after final failure. Clearly, the transverse cracks seen at lower loads have extended to form delaminations that proceed between the layers, in itiating ult imate failure.  The stress versus strain and photomicrographs for the specimens containing aramid pulp resin reinforcement in the ply drop region followed similar trends, namely, that a slight ju mp was observed in the stress-strain plots, at a value confirmed by the edge micrographs.For brev ity, further stress-strain plots and photomicrographs are not shown, but summarized in Tables 1 and 2.As shown in Table 1, for the of[0/ 0*/0] T laminate containing aramid pulp additive, the initial strain ju mp d id not occur until 757 M Pa.This value was confirmed (actually, bracketed) by the micrographs, which did not show any transverse resin cracks when a 322 MPa stress was reached, but a crack did appear after a load of 778 MPa was reached.Therefore, when co mpared to the control (epo xy-only resin), the laminate with aramid pulp in the region of the resin pocket had an in itial failure (t ransverse crack in the ply drop resin pocket) that was 2.75 times greater.The value of the ultimate failu re, also shown in Table  1, did not change significantly: fro m 924 MPa (control) to 966 M Pa (aramid pulp), an increase of less than 5%, and within the scatter of the test data. The results for the specimens containing PR-25 carbon nanofiber resin reinforcement in the ply drop region is also shown in Table 1.Note fro m Table 1 that the initial strain ju mp occurred at 265 MPa, similar to the control case.This was confirmed by the micrographs, which did not show a transverse resin crack at zero load, but one existed at 278 MPa.Therefore, the PR-25 carbon nanotube showed a slight (3.5%) drop in initial cracking when compared to the control case.In terms of ult imate strength, the the PR-25 showed a 9.1% decrease when co mpared to the control resin.
One explanation for the observed results in the[0/0*/0] T specimens is that the aramid pulp, due to its high ductility, may have increased the strain to failu re in the resin pocket, and hence delayed the initiat ion of the resin crack, since the strain level was controlled by the stiff adjacent 0° bounding plies.The addition of the carbon nanofiber did not provide for this increased ductility, so produced results similar to the control case.In both cases, ultimate failure was controlled by the failure of the bounding 0° plies, so occurred at similar stress values.

Quasi-static tests of of[+45/-45/0*/-45/+45] T
The results from the[+45/-45/ 0*/-45/+45] T laminates are shown in Table 2.These adjacent ply orientations were chosen to have a 'softer' longitudinal constraint when compared to the all-zero bounding plies considered previously.The stress-strain curves (not shown) are similar to those discussed previously, except for a slightly non-linear nature, common in laminates containing such a high percentage of 45° plies.In addition, the strain ju mp indicating transverse cracking was even more pronounced than in the all-zero degree laminates.As shown in Table 2, the epoxy-only control case shows a clear strain ju mp at 94.5 MPa, which is consistent with the photomicrographs showing no cracking at 88 M Pa, and developed transverse cracks at 123 MPa.
The stress-strain results for the[+45/-45/ 0*/-45/+45] T laminate, with the resin at the ply drop containing aramid pulp, indicated an init ial transverse crack at 106 MPa, which was consistent with the photomicrographs (no damage seen at 80 and 100 M Pa).The u ltimate failure of these laminates containing aramid pulp occurred at an average of 124 MPa, which is very similar to the control laminates (123 M Pa).
The stress-strain results for the[+45/-45/ 0*/-45/+45] T laminate, with the resin at the ply drop containing PR-25 carbon nanofiber, indicated an average nominal stress of 125 MPa for the appearance of the first transverse crack.This was somewhat (32%) higher value than the control laminates. Note the strain ju mp at 125 M Pa was consistent with the photomicrographs (no damage was visible at 103 MPa, transverse cracking was present at 130 MPa, and developed into delaminations at 154 MPa).The u ltimate failure of 160 MPa was higher than the control value of 123 MPa, an increase of 30%.
An exp lanation for the results is that the axial (loading direction) modulus and strength of the bounding plies were low (since all 45°), so the stiffness and strength of the reinforced resin played a larger ro le (than strain to failure) in the resin pocket failure.The addition of the aramid did not alter the strength of the resin pocket nearly as much as the strength increase due to the addition of the vapor grown carbon fiber, hence the minimal transverse crack and failu re strength increases seen with the aramid pulp, and the larger increases seen with the carbon nanofiber. The[0/0*/0] T laminates were tested under cyclic loading conditions:2 Hz at R = σ min /σ max = 0.1.The "S-N" results are shown in Figure 6.A lthough the scatter in fatigue is generally large (and no exception here), a classic power-law fit (linear in log-log p lot) of the data appears reasonable.The power law fit was chosen due to its long-standing use in the co mposite materials research field [12]. The pro ject funding level precluded additional test data, however, some overall fatigue behaviour can be identified. For examp le, the dramat ic reduction in fatigue strength for these thin laminates is apparent.There appears to be no clear distinct advantage or disadvantage to either of the nano-scale additives, when compared to the control epo xy-only case.A slight advantage may go to the aramid pulp reinforced resin, wh ich is consistent with the delay in the onset of the first transverse crack seen in the static results.

Fatigue of[45/-45/0*/-45/45]T
The fatigue results for the[45/ -45/0*/-45/ 45] T laminates are shown in Figure 7.Again, the rapid drop in stress versus life is apparent.Also, as before the power law relationship of the data is reasonable.As with the static data for this laminate, the best performing materials seem to be the epo xy reinforced with vapor gro wn carbon fiber.Consistent with the static results is the fact that the overall magnitude of the stress level is much lower than that observed in the all-0° laminates.

Stress Analysis Modeling
The addition of nano-scale additives to the resin may alter both the stiffness and fracture toughness of a co mposite laminate.Therefore, an understanding and quantification of these effects are important to assess potential benefits of the approach.The existence of ply drops in laminates creates complex stress states that contribute to the initiation and propagation of damage in laminates.The effects of cracks in the vicinity of the ply drop/resin pocket reg ion were modeled using 2D fin ite element analysis with the virtual crack closure technique (VCCT).Cracks were assumed to exist in one of three possible locations of a[0/0*/0] T laminate, as shown in Figure 8.Crack case #1 was assumed to exist as a delamination between the dropped ply and the bounding ply, just at the end of the dropped ply.Crack case #2 was assumed to exist as a delamination between the bounding plies, just at the location of convergence of the bounding plies at the end of the resin pocket.Finally, crack case #3 was assumed to exist as a transverse crack located at the end of the dropped ply.Normalized strain energy release rates (SERR) for Mode I (opening, G I ) and Mode II (shearing, G II ) were calculated for each condition.Mode III results are not shown here, since the tearing values for the specific cases studied herein were all zero.In this study, the value of the matrix modulus was varied fro m 2 to 5 GPa to investigate the effect of resin stiffness increase on the crack driving energy (epoxy, polyester, and vinylester resins with no reinforcement have, typically, moduli in the range of 3-4 GPa).  The results for crack case #1 are shown in Figure 9.The strain energy release rates are non-dimensionalized by mu ltip lying the G values by[(applied stress)2 * (ply thickness) / (longitudinal modulus of the ply)].In addition, the crack lengths are normalized by dividing the crack length by the thickness of a single ply.The results for this case clearly show that an increase in matrix modulus results in a significant decrease in G I and G II .In addition, as the crack length increases (shown as plots with a/t increasing fro m 0.25 to 0.75), the energy release rate G I increases, indicative of unstable (or likely progressing) crack growth.The value of G II as a function of crack length does not change significantly.Clearly, an increase in G Ic and/or the matrix modulus would improve the delamination resistance at this location.An improvement in G IIc would have a negligib le effect on delamination progression.
The results for crack location #2 are shown in Figure  10.At this location (delamination at the tip of the resin pocket), G II is zero.Note that an increase in the matrix modulus slightly reduces the opening crack driving force G I .Also, an increase in the crack length increases G I (unstable crack location).Note that the overall magn itude of SERR in Figure 10 is half of that observed in Figure 9 (crack location #1).An increase in G Ic would reduce the tendency for crack gro wth at this location.  Figure 11 depicts the crack growth potential of a crack in location #3 (transverse crack at the end of the dropped ply).Un like the previous two cases, increasing the matrix modulus produces an increase in G I .This may possibly account, somewhat, for the lack o f improvement seen in the[0/0*/0] T laminates containing the high modulus vapor grown carbon fiber.Note that with increasing crack length, the value of G I increases, meaning crack gro wth would be unstable.The notably higher value of SERR at location #3 (a tenfold increase with respect to crack location #1 and a twenty-fold increase over crack location #2) indicates that crack location #3 is the most likely site o f damage init iation (which is verified by the photomicrographs).

Conclusions
The purpose of this study was to investigate the damage development due to ply drops in tapered fiberglass laminates under tensile static and fatigue loading, and to assess the t effects of using nano-scale matrix reinforcement in the resin pockets that form at the end of the internally dropped plies.Laminate panels were fabricated with two types of lay-ups:[0/0*/ 0] T and[+45/-45/0*/-45/+45] T , where the * indicates the ply that was dropped.The laminates were fabricated using resin film infusion of Epon 862 epo xy resin. These films contained no nanofiber, 10 g/ m 2 of aramid pulp, and 10 g/m 2 of Applied Sciences, Inc carbon nanofiber PR-25-PS-OX.The fiberglass used was Owens Corning WindStrand Uniweave Knytex.
In the statically loaded specimens, the placement of an axial strain gage was useful in determin ing the stress level at the initiat ion of the first transverse crack.The utility of this method was verified using micrographs of the polished laminate edges.
The[0/0*/0] T laminates with aramid pulp had an initial failure (transverse crack in the ply drop resin pocket) stress that was 2.75 times greater than the epo xy-only control.The PR-25 carbon nanotube showed a slight 3.5% drop in init ial cracking stress when compared to the control case.In terms of ultimate strength, the aramid pulp specimens showed a 4.5% increase, whereas the PR-25 showed a 9% decrease when compared to the control resin.The[45/-45/ 0*/-45/45] T laminates with aramid pulp had an init ial failure (t ransverse crack in the ply drop resin pocket) that was 11.6% g reater than the epoxy-only control.The PR-25 carbon nanotube showed a 30% increase in initial cracking when compared to the control case.In terms of ultimate strength, the aramid pulp specimens showed a 1% increase, whereas the PR-25 showed a 30% increase when compared to the control resin.
The laminates were tested under cyclic loading conditions: 2 Hz at R = 0.1.A power-law fit o f the S-N data appeared reasonable.Under fatigue loading, there appeared to be a very slight advantage for aramid pulp and carbon nanofibers in the[0/0*/0]T and[45/-45/0*/-45/45] T laminates, respectively, consistent with the static results.The dramat ic reduction in fatigue strength for these thin laminates was apparent.
Fin ite element stress analysis modeling of the crack growth indicated the effect of having increased matrix modulus in the resin pocket zone, as well as increased value of laminate crit ical energy release rate.The modeling correctly indicated the sequence of damage initiating fro m the transverse cracks in the resin pocket ad jacent to the dropped plies.