3.1. Effects of a Hardness Gradient at Interfaces on Fatigue Life and Crack Propagation in LMCs
The AA1050/AA5754 N2 LMC systems were investigated in order to study the effects of integrating thin layers of a stronger material embedded in a softer matrix on the fatigue life and crack propagation. The local mechanical properties of the different layer materials after N2 processing can be seen in
Table 2.
The hardness was calculated as a mean value of the individual AA1050 and AA5754 layers, respectively, from both the L1/4 and L1/8 architecture. The hardness of the AA1050 layers is about 0.8 GPa and of the AA5754 layers about 1.4 GPa. Consequently, the hardness gradient at the laminate interfaces can be calculated to be about 0.6 GPa. However, it must be noted, that this estimation of hardness gradient at laminate interfaces does not take into account the formation of an interface affected zone at the immediate vicinity of an interface caused by different shearing behavior of dissimilar materials [
14,
68,
69]. This must also be considered for the nanoindentation results presented in the following sections.
The fatigue life diagrams (
S-N curves) of AA1050/AA5754 N2 laminate systems with two different architectures as well as of the constituent monolithic AA1050 N2 and AA5754 N2 materials are plotted in
Figure 4a. For each specimen tested, the number of cycles to crack initiation (
Ni) as well as the numbers of cycles to failure (
Nf) are shown. The determination of the number of cycles to crack initiation using a 0.1 Hz drop criterion of resonance testing frequency can be seen in
Figure 4c for a monolithic AA1050 sample and a AA1050/AA5754 laminated composite.
As can be seen in
Figure 4a, the fatigue life of a soft AA1050 matrix material can be enhanced significantly by integrating thin and strong interlayers of AA5754 forming a laminated structure. This can be observed in both the LCF and the HCF regimes. The fatigue life of the L1/4 architecture is lower for both regimes compared to the L1/8 architecture, where the harder AA5754 layer is positioned more towards the surface (see
Figure 4b). The increased fatigue life for the L1/8 architecture might be caused by effects of the interface affected zone [
14,
68,
69], different shear strain distribution in the respective laminate surface layers during ARB processing [
70] or a complex internal stress state upon cyclic loading due to the co-deformation of different aluminium layers with dissimilar hardness [
71]. The transition between the LCF and HCF regime is correlated to a threshold value which depends on the material of the layer at the surface [
16,
32,
46]. Below the threshold, in the HCF regime, the fatigue life of monolithic materials and laminated composites is primarily determined by the fatigue crack initiation stage: The numbers of cycles to crack initiation account for 95% to 99% of the numbers of cycles to failure for both monolithic materials and laminates. Above the threshold, in the LCF regime, the role of the crack propagation stage of fatigue life is significantly enhanced for the laminated composites compared to monolithic materials. The crack propagation stage accounts for 10% of numbers of cycles to failure for the AA1050/AA5754 N2 L1/4 composite and 13% for the AA1050/AA5754 N2 L1/8 composite compared to about 5 to 6% for the monolithic AA1050 N2 and AA5754 N2 materials, respectively.
Figure 4c shows a comparison of the drop of the resonance testing frequency between a laminated AA1050/AA5754 N2 L/8 specimen and a monolithic AA1050 N2 specimen tested in the LCF regime at 150 MPa (i.e., around 20 MPa above the threshold value). In both samples, the initiation of the macro-crack occurs at the respective surface layers consisting of the softer AA1050 material. As visible in the diagram, the resonance frequency of the laminated structure initially drops significantly slower with respect to the normalized numbers of cycles to failure
N/Nf as compared to the monolithic sample. This behavior can be attributed to a retardation of crack propagation at the vicinity of the interface to the harder AA5754 layer in the laminated composite [
46]. Moreover, macro-crack propagation captures a significantly higher fraction of the total fatigue life in the LMCs than for the monolithic materials, as indicated by
Ni,LMC and
Ni,Mono in
Figure 4c.
3.2. Effects of a Hardness Gradient at Interfaces on Fatigue Crack Growth (FCG) in LMCs
The effects of a periodic variation of strong and soft layers in a laminated composite and the associated result of having a periodic variation of the local driving forces at crack tips near the interfaces were studied using fatigue crack growth experiments on a AA1050/AA5754 N3(Nxr) laminate architecture. The local hardness of the different layers in the laminated section of the LMC architecture are listed in
Table 3 as mean values of all AA1050 and AA5754 layers, respectively.
As shown in
Table 3, no significant difference in hardness between AA1050 and AA5754 layers in the N3(Nxr) processed laminate and the hardness in the respective N2(N1r) processed monolithic materials can be observed. As the laminate and monolithic materials were recrystallized prior to each roll bonding process, this indicates that the local mechanical properties of the respective laminate layers and their constituent monolithic material can be considered comparable. The AA5754 layers in the LMC exhibit a hardness of 1.35 GPa and will be denoted as the stronger/harder layers in the following in opposition to the softer AA1050 layers with a hardness of 0.73 GPa. The resulting hardness gradient at the interfaces in the laminated composite can be calculated to be about 0.6 GPa.
Fatigue crack growth rates from increasing, decreasing, and constant stress intensity range tests for monolithic AA1050 and AA5754 as well as for constant stress intensity range tests for the AA1050/AA5754 laminated composite in crack arrester orientation are shown in
Figure 5.
No significant differences between crack growth rates measured by increasing, decreasing, and constant stress intensity range tests were found for monolithic AA1050 and AA5754 specimen, respectively. This suggests that there are no effects on the crack growth rate from prior loading history for decreasing tests and from transient crack growth effects for increasing tests, since the normalized
K-gradients for each test were chosen appropriately. The crack growth rate, da/dN, in monolithic materials between 5 and 12.5 MPa√m for AA1050 and 6–16 MPa√m for AA5754 can be described with respect to the applied stress intensity range, ∆
K, by the well-known Paris power-law relationship [
72,
73]:
where
C and
m are the scaling constants, characterizing the crack growth behavior in Region II. The Paris equation exponents
m were found to be about
m = 3.4 and 4.5 for the monolithic AA1050 and AA5754 materials, respectively (see
Table 4). The crack growth rates of the monolithic AA5754 material are lower compared to the AA1050 alloy. At a stress intensity range below about 6 MPa√m, a deviation from Region II crack growth behavior can be observed for AA5754, indicating the transition towards near-threshold Region I fatigue crack growth. This suggests that the threshold ∆
Kth of the technically pure aluminium AA1050 is lower than that of the AA5754 alloy. The near threshold fatigue crack growth behavior of the materials was not addressed experimentally, as measurement of crack growth rates below 5 × 10
−7 mm/cycle were associated with long measurement times at a testing frequency of 3 Hz on the servohydraulic testing system. Additionally, measurements of crack growth rates above 3 × 10
−4 mm/cycle were associated with increasing effects of plasticity resulting in mixed mode failure of specimen from plastic collapse and fatigue crack growth. The measurement data above this threshold value was discarded, as the crack length measured by the crack propagation gauges was considered inaccurate.
Regarding the laminated composite, the crack growth behavior between 7.5 and 20 MPa√m can be divided into two different zones, which are denoted as Region II-1 and Region II-2 in the following. The Paris equation exponents
m of the laminated composite were calculated to be to be about
m = 0.9 and 5.7 for Region II-1 and Region II-2, respectively (
Table 4).
As shown in
Figure 5, the fatigue crack growth rate of the laminated AA1050/AA5754 composite is significantly reduced compared to crack growth rates in both constituent monolithic materials. In Region II-1, the fatigue crack growth rate in the laminated composite at a constant ∆
K of 7.5 MPa√m is about the same as in the monolithic AA5754 and about 20% of the crack growth rate in monolithic AA1050 material. For FCG tests at higher constant ∆
K levels in Region II-1, further deviation of crack growth rates between the LMC and monolithic constituent material can be observed, as indicated by the lower exponent
m for the laminated composite. At ∆
K = 14.5 MPa√m, the fatigue crack growth rate behavior changes and can be described at subsequent higher stress intensity ranges using a different exponent
m for Region II-2. FCG tests at the transition between Regions II-1 and II-2 indicate the biggest decrease of fatigue crack growth rate in the laminated composite compared to the constituent materials (
Figure 5). The FCG rate of the LMC at 14.5 MPa√m was measured to be about 7% of the crack growth rate in monolithic AA5754 material.
A reduction of the fatigue crack growth rate in a laminated composite tested in crack arrester orientation of this magnitude compared to the constituent monolithic materials can only be explained by crack growth retardation effects associated with toughening mechanisms. In LMCs, these toughening mechanisms emerge from interactions between the process zone ahead of the crack tip and the periodic variation of mechanical properties at the interfaces due to the laminated architecture [
27].
The surface crack networks of the LMC specimen were investigated after testing in order to assess the crack growth behavior perpendicular to the interfaces of the laminated composites.
Figure 6 shows the fatigue crack growth paths in the laminated composite structure for experiments conducted at constant stress intensity ranges in Region II-1 (7.5 MPa√m to 12.5 MPa√m) and Region II-2 (15 MPa√m to 20 MPa√m), respectively.
At a low stress intensity range of 7.5 MPa√m, the crack path is orientated perpendicular to the layers and interfaces without being deflected significantly throughout the laminated structure. For increasing stress intensity ranges starting at 10 MPa√m, the crack path across the laminated section is increasingly impeded at the vicinity of interfaces. For experiments at constant stress intensity ranges of 12.5 MPa√m and above, a clear distinction of different toughening mechanisms at the interfaces can be observed: (a) crack deflection (orange arrows,
Figure 6) can be observed when the crack approaches interfaces from the softer (AA1050) layers towards the harder (AA5754) layers and (b) crack bifurcation (green arrows,
Figure 6) can be found when the crack approaches the interfaces from the harder towards the softer layers.
The crack deflection mechanism at the interface results in crack growth along the interface. As the experiments were conducted using the SE(B)-specimen geometry, the maximum bending stress is located at the symmetry plane above the notch root and is reduced along the span S. This limits the crack growth along the interfaces, as the driving force is reduced gradually. The deflection of the crack path at the interfaces, when cracks approach interfaces from the soft layers towards the harder layers, and consequent crack growth along the interface imply, that further crack extension along the loading axis requires a new crack to be nucleated in the adjacent layer. As this crack re-nucleation phase depends on the local (micro-) structural characteristics and the tests were operated under a constant far-field stress intensity range, it is evident that the toughening mechanism of crack deflection strongly promotes the damage tolerant fatigue behavior of the LMC.
Bifurcation of the cracks can be observed for the opposite case, where a crack approaches the interface from the harder towards the softer layer. This mechanism leads to a branching of the original crack front into two new separate cracks. The bifurcation angles are found to be around 45° in relation to the symmetry plane above the notch root, where the bending stress is at a maximum. Bifurcation of the crack front leads to local redistribution of the far-field crack driving force that is remotely applied using a constant stress intensity range ∆
K, as it is reallocated across multiple crack tips. Crack growth across multiple crack fronts reduces the overall rate of fatigue crack growth and thus further enhances the damage tolerant fatigue properties of the laminated composite as seen in
Figure 5.
Further investigation into these toughening mechanisms was done using synchrotron X-ray computed microtomography (SXCT), as these mechanisms described above were identified analyzing the crack networks at the surfaces of the laminated composites post-mortem. Results of the SXCT experiment on the AA1050/AA5754 N3(Nxr) laminated composite specimen fatigued at a constant stress intensity range of ∆
K = 17.5 MPa√m, as can be seen in
Figure 7.
The 3D-reconstruction of the crack network in the laminated specimen (
Figure 7a) shows that the crack network can be divided into a primary (green) and a secondary (blue) crack network. The primary crack network starts at the notch root and subsequently grows along the symmetry plane of the laminated SE(B)-specimen. At this symmetry plane, the bending stress is at a maximum and promotes the highest tensile stresses perpendicular to the direction of the crack path and thus the highest driving forces at the crack tip. At the vicinity of the interfaces, the crack path is deflected due to interaction mechanisms between the process zone ahead of the crack tip and the variation of local mechanical properties at the interfaces. This results in the formation of the secondary crack network associated with the prevalent toughening mechanisms of crack deflection and crack bifurcation.
The crack networks in four cross sections of the 3D-tomogram at different positions
z = 0.05
B,
z = 0.33
B,
z = 0.66
B and
z = 0.95
B across the thickness
B of the LMC specimen were analyzed in
Figure 7b–d, respectively. Crack deflection (orange arrows) and crack bifurcation mechanisms (green arrows) can be conclusively identified in cross sections of the crack network both near the surfaces (
Figure 7b,e) as well as in the volume (
Figure 7c,d) of the laminated composite sample.
These findings emphasize the magnitude of the effects on damage tolerant fatigue crack growth behavior in LMCs that toughening mechanisms at interfaces in laminated metallic composites with dissimilar hardness can produce.
3.3. Effects of a Combined Gradient of Hardness and Elastic Properties at Interfaces on Fatigue Life and Crack Propagation in LMCs
Investigations into the influence of a combined difference in strength and elastic properties at interfaces between layers on fatigue life and the role of toughening mechanisms at interfaces on crack propagation was studied on a AA7075/DC05 N2 L1/4 laminated composite architecture. The local mechanical properties in terms of hardness and Young’s modulus of the different layer materials after N2 processing as well as the resulting gradients at interfaces can be seen in
Table 5.
The average hardness was measured to be 1.55 GPa in the AA7075 aluminium layers and 2.66 GPa in the DC05 steel layers of the laminated composite. Consequently, the hardness gradient ∆
H at the laminate interfaces can be calculated to be about 1.1 GPa. The Young’s moduli of the respective layers were assumed to be 70 GPa for AA7075 and 210 GPa for DC05 based on literature data [
74,
75], resulting in a gradient in elastic modulus ∆
E of 140 GPa at the AA7075/DC05 interfaces.
Figure 8a plots the fatigue life diagrams (
S-N curves) of the AA7075/DC05 N2 L1/4 laminate architecture as well as of the constituent monolithic AA7075 N2 and DC05 N2 materials. Again, the number of cycles to crack initiation (
Ni) as well as the numbers of cycles to failure (
Nf) are shown for each fatigue test. The percentile of the fatigue crack propagation phase on the total fatigue life is specified for the LCF regimes of the laminate and monolithic materials, respectively. The S-N diagram reveals that the fatigue life of the AA7075 N2 matrix material can be enhanced in the LCF regime as well as in the HCF regime by integration of a thin DC05 steel layer.
Previous studies by the authors [
15,
16] on different laminated Al/Steel composite systems with the same architecture as in
Figure 8c revealed that the enhancement of the endurable maximum stress amplitudes ∆
σmax/2 in the HCF regime correlates with the reduction of the maximum bending stress in the outer aluminium layer due to load transfer associated with the higher elastic modulus of the steel layer. FEM analysis of the stress distribution upon loading bending of Al/Steel laminated beams showed that the introduction of the steel layer near the surface of the laminate reduces the tensile stress at the aluminium surface layer. Using the FEM model presented by Kümmel et al. [
15], the tensile stress at the outer AA7075 layer can be calculated to be reduced by 19% for the L1/4 architecture (
Figure 8c) compared to monolithic AA7075 material. The fatigue limits of the AA7075/DC05 N2 L1/4 composite and the AA7075 N2 monolithic material can be estimated at 210 MPa and 170 MPa, respectively. Calculating an effective fatigue limit of the AA7075/DC05 composite and taking into account the FEM-calculated 19% reduction of the maximum tensile stress in the outer AA7075 layer resulting from the load transfer into the adjacent DC05 steel layer, the resulting effective endurable maximum stress amplitude ∆
σmax,corr./2 amounts to 170.1 MPa. This fits perfectly to the obtained fatigue limit of the monolithic AA7075 N2 material of 170 MPa. The numbers of cycles, which are consumed to initiate a crack in the HCF regime, amount to 96–99% of
Nf for both the laminated composite and constituent monolithic materials. Again, it becomes evident that in the HCF regime the fatigue life of the LMC is primarily determined by the fatigue crack initiation phase, as is the case for the monolithic AA7075 and DC05 materials and the overall role of crack propagation processes in the LMCs subjected to loadings in the HCF regime, which is relatively small.
This load transfer effect enhances the fatigue life of the LMC in the LCF regime as well, compared to the monolithic AA7075 material. Additionally, as reported in
Section 3.1. for the Al/Al LMC systems, in the LCF regime, the role of fatigue crack propagation is promoted in the AA7075/DC05 LMC system compared to the constituent monolithic materials. As indicated in
Figure 8a, the crack propagation phase accounts for 17% of the numbers of cycle to failure for the AA7075/DC05 composite compared to about 5 to 7% for the monolithic AA7075 N2 and DC05 N2 materials. This increased percentile of the crack propagation phase on the overall fatigue life in the laminated composite must be associated with toughening mechanisms impeding the propagating crack front at the vicinity of the interfaces.
Figure 8b shows SEM images of the fatigue crack propagation paths through the laminated structure of the AA7075/DC05 N2 L1/4 architecture of specimens fatigued in the LCF and HCF regime, respectively. A clear difference in the crack propagation behavior in the LMC architecture between the LCF regime compared to the HCF regime can be seen. In the HCF regime, the crack path propagates relatively straight through both the interfaces. A slight deviation in crack path trajectory away from the initial orientation perpendicular to the interfaces can be observed when the crack approaches the interface towards the harder and stiffer steel layer. Before penetrating the steel layer, the crack propagates about 100 µm along the first interface. In the LCF regime, a more distinct crack network can be observed. Crack propagation in 45° orientation towards the surface can be observed in the outer AA7075 layer. Some degree of necking occurring in the DC05 steel layer hints to promoted activities associated with plasticity involved in the crack propagation process at this stage. Applying only a post-mortem analysis of the crack networks in the laminated AA7075/DC05 composites does not deliver sufficient evidence to identify individual toughening mechanisms associated with the enhanced role of fatigue crack propagation in the LCF regime.
In order to gather a better understanding of the prevalent toughening mechanisms obstructing fatigue crack propagation at the vicinity of interfaces in AA7075/DC05 N2 L1/4 laminated composites, in-situ LCF and HCF fatigue experiments have been conducted inside the large chamber SEM.
The experiment in the LCF regime was conducted at a maximum bending stress amplitude ∆
σmax/2 of 295 MPa as shown in
Figure 9a.
Figure 9b shows the crack network after 30.8k loading cycles. Multiple small cracks can be observed at the front surface. The formation of these small cracks in a AA7075 alloy is known to be associated with the brittle behavior of small intermetallic particles (e.g., Mg
2Si and Al
7Cu
2Fe) [
76]. Two larger cracks have been formed at the bottom surface of the specimen at this stage. The larger crack on the right side was nucleated at the surface aluminium layer below the loading anvil. At this position, the tensile stress at the surface is at a maximum, resulting from the cyclic loading in three-point bending mode. This crack has propagated through half of the outer AA7075 layer at an angle of about 45° with respect to the orientation of the interface. From this position on, the crack path trajectory changes and the crack propagates perpendicular to the interfaces (
Figure 9c). As the crack approaches the immediate vicinity of the interface towards the steel layer after 31.5k loading cycles, the crack path gets deflected prior to the crack tip reaching the interface (
Figure 9c). Upon reaching the aluminium/steel interface, crack propagation along the interface of about 350 µm was observed during the subsequent sequence of approximately 1k loading cycles (
Figure 9d).
During this period, the formation of interface delamination cracks associated with stress redistribution effects at the opposing steel/aluminium interface was observed. After 32.8k loading cycles, further crack propagation into the inner AA7075 layer of the LMC architecture starting at one of the delamination cracks was observed (
Figure 9e) before the original crack begins to penetrate the steel layer. This crack bridging mechanisms of the steel layer leads to the formation of a complex crack network with simultaneous crack propagation occurring in the aluminium and steel layers (
Figure 9f, 33k loading cycles). The crack propagation in the DC05 steel layer is associated with visible plastic activity around the crack tip. Although no necking of the steel layer in the in-situ LCF experiment was found, as observed in
Figure 8b, the plasticity is indicated by the formation of shear bands ahead of the crack tip at an angle of about 45° to 60° in relation to the direction of loading in
Figure 9f. Upon coalition of the two individual cracks, the test was stopped, as the crack network reached the bonding plane of the final N2 ARB processing step (
Figure 9a).
A second in-situ fatigue experiment was conducted in the HCF regime at a maximum bending stress amplitude ∆
σmax/2 of 210 MPa, as shown in
Figure 10a. Prior to the in-situ experiment in the large chamber SEM, the specimen was pre-fatigued externally on a vibrophore testing machine at the same maximum bending stress amplitude of 210 MPa for 9.1 Mio cycles. The initiated crack after 9.1 Mio loading cycles using the method described above can be seen in
Figure 10b. The crack was nucleated at the position of the maximum tensile stress in the outer AA7075 layer resulting from the cyclic three-point bending loading and propagates at a perpendicular orientation towards the interfaces. As in the LCF experiment, multiple small cracks originating at intermetallic particles can be seen on the front surface. No major contributions of these cracks on the propagation of the major crack were observed during the in-situ experiments. After 20 k additional loading cycles, the crack reaches the aluminium/steel interface (
Figure 10c). As in the LCF regime experiment, a deflection of the crack path can be observed before the crack tip reaches the interface. Consequently, the crack approaches the interface at a high deflection angle and propagates along the interface for about 40 k loading cycles, as can be seen in
Figure 10d.
At this stage, multiple events associated with local plasticity can be observed at different positions in the DC05 steel layer at the vicinity of the interface, indicated by secondary electrons contrast edge effects at roughened surfaces. Consequently, at one of these positions, a crack is nucleated into the adjacent steel layer. No formation of interface delamination cracks at the opposing steel/aluminium interface associated with stress redistribution effects was observed during this in-situ experiment. As can be seen in
Figure 10e, after an additional 20 k cycles, the crack propagated through the steel layer at an angle, returning to the symmetry plane where the driving force on crack propagation is at a maximum due to the cyclic loading condition under three-point bending mode. After reaching the steel/aluminium layer, the crack deflects and propagates along the interface before penetrating the inner AA7075 layer (
Figure 10f). This crack deflection at the steel/aluminium interface is much less pronounced when the length of crack that propagates along the interface (about 100 µm for the steel/aluminium interface vs. 250 µm for the aluminium/steel interface) is regarded. The less pronounced crack deflection is also visible by comparing the loading cycles for the crack propagation along the interfaces (below 5 k for the steel/aluminium interface vs. 60 k for the aluminium/steel interface). As for the loading cycle comparison, it must be stated again that the experiments were conducted under force control. The experiment was stopped once the crack reached the bonding plane of the final N2 ARB processing step. The crack network of this in-situ fatigued HCF specimen post-experiment, as can be seen in
Figure 10a, resembles the crack network of the HCF fatigue-life specimen shown in
Figure 8b, indicating an appropriate representation of HCF crack propagation using the described experimental procedure.
Using in-situ fatigue experiments, different toughening mechanisms associated with the interaction of the process zone ahead of the fatigue crack and the interfaces in laminated metallic composites with a combined gradient in hardness and elastic properties could be identified. The influence of these individual mechanisms was assessed in a semiquantitative manner. The findings can explain the enhanced fatigue crack propagation behavior of the laminated composites in the LCF regime associated with the prevalent toughening mechanisms.
3.4. Effects of the LMC Architecture Combined with Gradients of Hardness and Elastic Properties at Interfaces on Fatigue Life and Crack Propagation
Previous studies by the authors [
15] revealed, that fatigue properties of LMCs can be specifically tailored by modifying the laminate architecture accordingly. In order to find optimized laminate architecture designs regarding fatigue life and crack propagation, laminates utilizing different combinations of material inhomogeneity effects in terms of gradients of hardness and elastic properties at interfaces were tested. The local mechanical properties in terms of hardness and Young’s modulus of the different AA2024, Ti-G1 and DC05 layers as well as the resulting gradients at interfaces can be found in
Table 6.
Based on the respective hardness and elastic properties of the individual layers (
Table 6), the highest gradients ∆
H and ∆
E in the laminated composite systems can be found at the AA2024/DC05 interfaces and the smallest gradients ∆
H and ∆
E at the AA2024/Ti-G1 interfaces, with the respective gradients of hardness and elastic properties of the Ti-G1/DC05 interfaces ranging in between.
Figure 11a shows the architectures of the different tri-material laminated composites investigated, containing thin layers of AA2024, Ti-G1 and DC05 near the respective surfaces (shell structure) as well as a AA2024 material core structure. The near surface AA2024, Ti-G1, and DC05 layers in the shell structure of the laminate have about the same thickness in average in all three laminate architectures, leading to the same nominal density for all LMCs.
The fatigue life diagrams (
S-N curves) of the Ti-G1/DC05/AA2024, Ti-G1/AA2024/DC05, and AA2024/Ti-G1/DC05 N3 laminate architectures as well as of the constituent AA2024 N3 and Ti-G1 N3 materials are plotted in
Figure 12a. In all three laminated composite systems, the fatigue life in both the LCF as well as HCF regime can be improved significantly by the integration of thin titanium and steel layers compared to the monolithic AA2024 N3 material. A comparison of the
S-N curves at the respective LCF to HCF region transition areas reveals a gradual enhancement in terms of endurable maximum stress amplitude and numbers of cycles to failure, indicating increasing levels of optimized laminate architecture designs regarding fatigue life properties.
The best resistance against crack initiation for the different LMC systems could be achieved in laminates where the titanium layers are positioned at the respective surfaces: Ti-G1/DC05/AA2024 and Ti-G1/AA2024/DC05. Comparing these two laminate architectures, the resistance against crack initiation is enhanced for the Ti-G1/DC05/AA2024 laminate architecture compared to the Ti-G1/AA2024/DC05 architecture. This can be explained again by the load transfer from the titanium surface layer into the adjacent DC05 steel layer due to the gradient in elastic modulus, thereby reducing the effective tensile stresses at the bottom surface titanium layer.
The best resistance against crack propagation was observed for the AA2024/Ti-G1/DC05 laminate architecture, where the crack propagation phase accounted for 20% of the numbers of cycle to failure in the LCF regime compared to 11% for the Ti-G1/AA2024/DC05 architecture, 10% for the Ti-G1/DC05/AA2024 architecture and about 6 to 7% for the monolithic Ti-G1 N3 and AA2024 N3 materials. This coincides with the findings presented in
Section 3.1 and
Section 3.3 and can be explained by a reduction of the local driving force at the crack tip, when the crack approaches the interfaces to the stronger and stiffer layers. For the AA2024/Ti-G1/DC05 architecture, this is the case for both the AA2024/Ti-G1 and the Ti-G1/DC05 interfaces (see
Table 6), where the crack approaches the interface from the layer denoted first to the one denoted second, respectively.
Figure 11b shows SEM images of the fatigue crack propagation paths through the laminated structure of the Ti-G1/DC05/AA2024 and Ti-G1/AA2024/DC05 LMCs fatigued in their respective LCF regimes.
In
Figure 12b, the
S-N curves are plotted in relation to the respective density ρ of the laminated composites and monolithic materials. The monolithic DC05 N3 steel exhibits by far the worst specific fatigue properties due to its high density. The specific fatigue life of the Ti-G1/DC05/AA2024 and Ti-G1/AA2024/DC05 LMC architectures are significantly improved compared to either of the constituent monolithic AA2024 N3, Ti-G1 N3, and DC05 N3 materials.
These findings clearly demonstrate that by an appropriate selection of the constituents, by adjusting their mechanical properties and by an intelligent design of the laminated metallic composite the fatigue life and cyclic crack propagation behavior of LMCs can be significantly improved.