Mode I fatigue crack growth induced by strain-aging in precipitation-hardened aluminum alloys

Theoretical and Applied Fracture Mechanics 104 (2019) 102340

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Mode I fatigue crack growth induced by strain-aging in precipitationhardened aluminum alloys Samsol Faizal Anisa, Motomichi Koyamab, Shigeru Hamadac, Hiroshi Noguchic,

T

⁎

a

Department of Mechanical Engineering, University of Technology Malaysia, Jalan Sultan Yahya Petra, 54100 Kuala Lumpur, Malaysia Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan c Department of Mechanical Engineering, Kyushu University, 744 Motooka, Nishi-ku, Fukuoka 819-0395, Japan b

A R T I C LE I N FO

A B S T R A C T

Keywords: Fatigue crack growth rate Aluminum alloys Fatigue striation Mode I crack Strain aging

In this study, the effects of Mg on the fatigue crack growth (FCG) characteristics of precipitation-hardened Al6061 alloys containing Zr and excess Mg are examined. The growth behavior of microstructurally large cracks is investigated via rotating-bending fatigue tests conducted at room temperature. Analyses of the crack propagation and striation features show that excess Mg promotes the occurrence of Mode I fatigue cracks. These facts suggest that dynamic strain aging due to Mg restricts dislocations motion, resulting in greater work hardening, and generates a large number of active slip systems, leading to more non-localized slip and a large area of striation formation. Consequently, it is concluded that this phenomenon induces highly stable crack growth in Al alloys with excess Mg, which influences the fatigue crack growth rate (FCGR) scatter of microstructurally large cracks. A reasonable mechanism of strain-aging-induced Mode I FCG in Al alloys is proposed on the basis of the morphological aspects of fatigue striation formation.

1. Introduction Automotive lightweight construction based on newly designed aluminum alloys is one of the most promising solutions for addressing the growing need to reduce fuel consumption and CO2-emission [1,2]. Precipitation-hardened Al alloys such as Al6061-T6 alloys are advantageous for the fabrication of automotive parts owing to their high corrosion resistance, lightweight, high strength, good weldability, and reasonable fatigue resistance [3,4]. However, there remains a scope for improving their fatigue limit and life. In particular, Al6061-T6 alloys do not have a distinct fatigue limit associated with fatigue crack nonpropagation [5] and show large scatter of their fatigue life and fatigue crack growth rate (FCGR) [6,7] compared to steels [8]. Both of the facts are critically important for engineering design. Hence, fatigue behavior of Al6061-T6-based alloys has been explored to endow a high-strength aluminum alloy with a distinct fatigue limit. In a previous study [9,10], a distinct fatigue limit and associated coaxing effect were reported to appear in an excess Mg-doped Al6061T6 alloy where Mg remained as a solute state even after precipitation of Mg2Si. Specifically, the excess solute Mg results in dynamic strain aging (DSA), assisting occurrence of fatigue crack non-propagation at the fatigue limit [9,10]. However, the excess Mg increased the grain size, which decreases the tensile strength. To solve this problem, a Al6061-

⁎

T6 alloy [11] containing excess Mg and Zr was recently developed. The simultaneous addition of Mg and Zr provided finer grains than those in the previous Al6061-Mg alloy [9,10], which realized high tensile strength with a distinct fatigue limit. Furthermore, FCGR of microstructurally large cracks was lower than that in an Al6061-Zr alloy without an excess Mg. As mentioned above, the disadvantages of Al6061-T6 alloys are (1) no distinct fatigue limit and (2) large scatter in FCGR and associated fatigue life. After the successful alloy design for distinct fatigue limit, we next aim to clarify the scatter characteristics of the newly developed Al6061-Mg-Zr alloy. In this context, it is necessary to clarify the effect of excess Mg on fatigue crack propagation to determine the underlying factors that cause the scatter. In general, the scatter in FCGR appears particularly in microstructurally small crack growth [12], in which microstructure effects are significant when the crack length increases over several times the grain size, the scatter in FCGR markedly decreases [13], which generally involves crack growth through activations of two alternative slip systems [14] at the crack tip (Mode I mechanism). However, Al6061-T6 alloys exhibit large scatter of their fatigue crack propagation life, even when the crack length exceeds several times the grain size, which seems to arise from ease of zigzag fatigue crack growth (Mode II mechanism) [7]. The significant scatter in the crack propagation life until the large crack size increases a degree

Corresponding author. E-mail address: [email protected] (H. Noguchi).

https://doi.org/10.1016/j.tafmec.2019.102340 Received 21 March 2019; Received in revised form 27 August 2019; Accepted 28 August 2019 Available online 29 August 2019 0167-8442/ © 2019 Elsevier Ltd. All rights reserved.

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Nomenclature AS AO c dl/dN E l

N R SR φ ν σy σ∞ ωp

striation area observation area half crack length crack growth rate Young’s modulus crack length

number of cycles stress ratio striation ratio crack tip opening displacement Poisson’s ratio yield strength remote stress size of plastic zone

Fig. 2(a)). The tensile tests were performed using a strain rate of 3 × 10−4 s−1 at room temperature. One specimen test for each Al alloys was used in order to find the mechanical properties of the Al alloys. Furthermore, rotating-bending fatigue tests were conducted using specimens with the geometry shown in Fig. 2(b). To eliminate the workhardened layer, all the specimens were mechanically polished and subsequently electropolished at 30 V and 323 K in a solution containing 2% gelatin, 1.2% oxalic acid dehydrate, and 96.8% phosphoric acid with a concentration of 85%. The fatigue tests were conducted at room temperature using an Ono-type rotating-bending fatigue test machine, with a stress ratio (R) of −1 and a frequency (f) of 60 Hz. Because this type machine was under the four-point bending, the central parts on 50 mm long region had a constant bending moment. The stress was defined as the nominal bending stress at the minimum section. To observe the fatigue crack growth and to measure the crack length, the fatigue test was stopped at predefined cycles. The fatigue crack growth was observed using a replica technique, in which plastic replicas were obtained in the unloaded condition. The replica technique is a method for crack detection, the detailed procedures of which are described elsewhere [25]. The crack length refers to the length along the circumferential direction of the specimen, and it was observed using a microscope. Fig. 3 shows a schematic illustration of the measurement of the crack length. The crack length measurement was classified into three categories, namely, crack initiation, crack propagation, and crack growth with branches. The crack length measurement data were used to plot a graph of crack length versus number of cycles. The secant method or point-to-point technique [26] was used to determine the crack growth rate dl/dN at a particular crack length using the following equation:

of the scatter in fatigue life of Al6061-T6 alloy. Therefore, we here note an excess Mg effect on the scatter of crack propagation at crack lengths over several grain sizes. More specifically, Mg-driven DSA is expected to affect the fatigue crack propagation behavior of the Mg-doped Al6061T6-based alloys. Based on this background, the present study aims to clarify the effect of excess Mg on the propagation behavior of relatively-large cracks in Al6061-T6-based alloys in terms of the scatter factor of FCGR. The effect of DSA on Mode I crack growth in Al6061-T6 alloys is investigated through characterizing fatigue striations on the fracture surface. The fractographic analysis on fatigue striation is not limited to the striation spacing, but also includes the topographic characteristics of the striations. The profile development of fatigue striations can be correlated with the behavior of Mode I crack growth, which thereby enables detailed discussion on propagation mechanisms of the relatively large cracks. Accordingly, we present mechanism-based interpretations of the probabilistic behavior of the crack growth in the highstrength aluminum alloys with and without excess Mg. 2. Experimental procedure 2.1. Material This study considered two newly developed precipitation-hardened Al6061-T6-based alloys [11], namely Al6061-Zr alloy and Al6061-ZrMg alloy, the chemical compositions of which are summarized in Table 1. The main characteristics of these two alloys that distinguish them from ordinary Al6061-T6 alloys are the addition of Zirconium (Zr) in both alloys and the presence of excess Magnesium (Mg) in the latter. Excess Mg is used to promote DSA [15–18], which has been reported to facilitate fatigue crack non-propagation in the infinite life regime. Accordingly, precipitation-hardened Al alloys containing excess Mg exhibit a distinct fatigue limit [9,10]. Furthermore, Zr is added to Al6061 alloys to improve their tensile strength by decreasing the grain size [19–22], as strength deterioration due to Mg addition is a drawback of Al6061-Mg alloy [9]. In the fabrication stage, billets of the alloys (diameter, ~155 mm) were prepared by semi-continuous casting. Then, homogenization treatment was applied to the billets at 823 K for 14,400 s. Therefore, the billets were extruded to form round bars (diameter, 23 mm) at 773 K. The extruded round bars were then solution-treated in an air furnace at 813 K for 3600 s and immediately water-quenched. Finally, T6 aging was performed [23,24] at 463 K for 14,400 s. Figs. 1(a) and (b) show the microstructures of the Al6061-Zr and Al6061-Zr-Mg alloys, respectively. Many defects were observed on the surfaces of the alloy specimens. Figs. 1(c) and (d) show the magnified images of the defects of the two Al alloys. The sizes of the defects in both the alloys were nearly the same (diameter, ~ 1–5 µm). The results of extreme value analysis are shown in Figs. 1(e) and (f); √areamax is the maximum size in the inspection area S0.

(l − li ) ⎡ dl ⎤ = i+1 , (Ni + 1 − Ni ) ⎣ dN ⎦at ⎛ li + 12+ li ⎞ ⎝

(1)

⎠

where l is the crack length and N is the number of cycles. 2.3. Fractographic analysis The fractographic features were observed using a scanning electron microscope (JEOL JSM IT300). The striation spacing was measured in six different regions, located at ~2000, 2500, 3000, 3500, 4000, and 4500 µm from the initiation site of the main crack. Fig. 4(a) shows a magnified image of the crack initiation site. The initiation point was identified by tracking the direction of the fatigue crack propagation. The presence of a gouge mark was one of indicators used to determine the crack initiation site [27]. The average striation spacing was then determined from around 20 striations measured in each selected region. In addition, the analysis of striation spacing is expanded using the Table 1 Chemical composition of the precipitation-hardened Al alloys (wt%).

2.2. Tensile and fatigue tests Plate specimens of the alloys (gauge dimensions: width, 4 mm; thickness, 1 mm; length, 30 mm) were used for tensile tests (see 2

Element

Si

Fe

Cu

Mn

Mg

Cr

Ti

Al

Zr

Al6061-Zralloy (wt%) Al6061-Zr-Mgalloy (wt%)

0.55 0.55

0.21 0.2

0.24 0.23

0.09 0.09

0.9 1.39

0.26 0.26

0.02 0.02

bal. bal.

0.16 0.14

Theoretical and Applied Fracture Mechanics 104 (2019) 102340

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y 8 F [%]

y 8 F [%]

99.95

(e) 6

99.9 99.8

4

99.5 99.0 98

2

0

95 90 80 50 10 1 0.1 0

99.95

(f) 6

99.9 99.8

4

99.5 99.0 98

2

95 90 80 50

0 S0 = 0.091 mm

2

4 6 areamax [μm]

10 1 0.1 0

2

8

S0 = 0.091 mm

2

4 6 areamax [μm]

2

8

Fig. 1. Microstructures of the precipitation-hardened Al alloys: (a) Al6061-Zr alloy and (b) Al6061-Zr-Mg alloy, (c) magnified image of Al6061-Zr alloy, (d) magnified image of Al6061-Zr-Mg alloy, (e) defect size statistics of Al6061-Zr alloy, and (f) defect size statistics of Al6061-Zr-Mg alloy.

To investigate the behavior of the Mode I crack, we evaluated the striation ratios for both Al alloys using measurements from seven different regions located between 500 and 3500 µm from the crack initiation site, as shown in Fig. 4(b). Fig. 4(c) shows the size of the examination area in each region. The striation ratio SR was calculated as follows:

central limit theorem to determine an approximate value of the population mean. It was done in the region of 4000 µm from the crack origin. The central limit theorem states that the distribution of the sample mean tends toward a normal distribution and its standard deviation shrinks when the sample size increases or the samples contain relatively large numbers of observations. In other words, based on the central limit theorem, the sample means become close to the population mean by analyzing a sufficient number of data. Hence, three hundred data with thirty sample sets which contain ten observations data of striations spacing in each Al alloy was examined. The sampling means were rounded up to four decimal places for analysis, and the normal probability plot were then used to assess the normality of the data set.

SR =

AS , AO

(2)

where AS is the striation area and AO is the observation area. The observation area was always normal to the electron beam position to prevent measurement errors in the striation spacing and striation area. 3

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(a)

(b)

Loading direction Fig. 2. Sample geometries (all units in mm): (a) tensile tests and (b) rotating-bending fatigue tests.

serrations in the stress-strain curve of Al6061-Zr-Mg alloy. The magnified image in Fig. 6(b) clearly shows the occurrence of serrations. Table 2 summarizes the differences between the mechanical properties of the two types of precipitation-hardened Al alloys.

In addition, the fatigue striation morphology of both sides of the fracture surface was investigated to determine whether it exhibits peakto-peak or peak-to-valley matching. The matching position was set by selecting two reference points having matching signs on both sides of the fracture surface, denoted by P1, P2 and V1, V2, as shown in Fig. 5(a)(i) and (ii), respectively. For instance, the peak point P1 in Fig. 5(a)(i) corresponds to the valley point V1 in Fig. 5(a)(ii). To reduce the error, two reference points should be selected from the other side of the fracture surface to confirm the exact location and to avoid inaccuracy due to rotation images. The matching area of fatigue striation was identified by measuring the distance between the two reference points and selected corner of a rectangular area. Fig. 5(b)(i) and (ii) show magnified images of the matching area of fatigue striation in Al6061-Zr-Mg alloy. The matching location was then determined by considering the same distance between a reference point and the selected point. Furthermore, the roughness profiles of the fracture surface were analyzed using 3D scanning electron microscope (SEM) images. Three SEM images of each Al alloy were obtained by tilting the specimen plane at +5°, 0°, and −5°, 4000 µm from the crack initiation site. The 3D fracture surface was then constructed automatically using an image processing program called MeX [28]. To analyze the surface roughness at the formation of striation, the roughness parameters, namely, the average roughness profile, Ra, and the average maximum height of the roughness profile, Rz [29], were examined using MeX, which is compliant with ISO 4287:1997 [28]. The implications of the results with regard to striation formation as well as the fatigue mechanism are discussed below.

3.2. Fatigue tests 3.2.1. S-N curves and fatigue limit characteristics Fig. 7 shows the graphs of stress amplitude versus number of cycles to failure (S-N) of the two types of precipitation-hardened Al alloys. In the finite life regime, Al6061-Zr alloy does not have a fatigue limit, which indicates the occurrence of failure, even after 107 cycles. By contrast, despite the additional Zr content, the fatigue strength of Al6061-Zr-Mg alloy is similar to that determined in previous studies [9,10], with a distinct fatigue limit. This phenomenon is attributed to the effect of strain aging, which strengthens the material around the front of the crack tip. Fig. 8 shows three non-propagating cracks at a stress amplitude of 143 MPa. The above-mentioned findings reflect the ability of Al6061-Zr-Mg alloy to resist the growth of a small crack. In addition, it was observed that the fatigue cracks easily coalesced at high stress amplitudes of ~250 MPa owing to the ease of crack initiation. Moreover, the fatigue cracks in Al6061-Zr-Mg alloy failure earlier than expected at stress amplitude of 160 MPa and 180 MPa owing to the crack coalescence. However, many non-propagating cracks occurred at lower stress amplitudes, especially in Al6061-Zr-Mg alloy, owing to the strengthening mechanism in front of the crack tip. Hence, to examine the FCGR with the purpose of minimizing crack coalescence and non-propagating cracks, we focused on a stress amplitude of 200 MPa and investigated the propagation-mode characteristics. At this stress amplitude, the main crack was identified as a single crack growth, which did not coalesce with other cracks to form a dominant crack.

3. Experimental results 3.1. Tensile tests

3.2.2. Fatigue crack growth behavior on specimen surface In the present study, fatigue crack behavior consists of two periods:

Fig. 6(a) shows the results of the tensile tests. As can be seen, Al6061-Zr alloy has a smooth stress-strain curve. By contrast, there are

Fig. 3. Schematic illustrations of crack length measurement: (a) crack initiation at early stage, (b) crack initiation and propagation, and (c) crack propagation with branches. 4

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Fig. 4. SEM images of the fatigue fracture surfaces: (a) magnified image of the fatigue crack initiation site, (b) seven regions of fatigue striation evaluation, and (c) size of the observation area for each region.

Fig. 6. Tensile test results of the precipitation-hardened Al alloys: (a) stressstrain curve and (b) magnified image of the serration. Fig. 5. SEM images of the matching area in Al6061-Zr-Mg alloy: (a) viewed from both sides of the fracture surface at 3500 μm from the crack origin and (b) magnified images of fatigue striation.

Table 2 Mechanical properties of the precipitation-hardened Al alloys.

microstructurally small crack and microstructurally large crack. The total fatigue life is influenced by crack initiation, interaction of multiple cracks, and the propagation mode of the cracks. Fatigue cracks were initiated at defects on the specimen surface owing to the high stress concentrations induced by the defects. By considering the crack growth without crack interaction, the fatigue cracks in Al6061-Zr-Mg alloy were generally initiated earlier than those in Al6061-Zr alloy but they grew more slowly after 2 × 105 cycles, particularly for the main crack as shown in Fig. 9 [11]. The mode I crack growth is considered when the stable crack growth occurs, and the comparison of the main crack growth behavior in both Al alloys can be observed from Fig. 10. To investigate the Mode I fatigue crack growth behavior in greater detail, we observed the fracture surfaces using an SEM.

Alloys

0.2% proof strength (MPa)

Tensile strength (MPa)

Elongation (%)

Al6061-Zr alloy Al6061-Zr-Mg alloy

327 294

350 332

7.38 8.66

3.3. Fractographic investigation The growth of a microstructurally large fatigue crack involves alternating plastic deformations along two symmetrical slip systems. Through fractographic observations, Mode I fatigue crack growth could be examined on the basis of the formation of fatigue striations on the fracture surfaces as well as the well-known mechanism for describing striation formation, namely successive blunting and resharpening of the crack tips [30,31]. Nevertheless, the development of the fatigue striation profile varies with the type of material and is affected by cyclic 5

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previous study in which specimens with drilled holes were employed [35]. Furthermore, by considering thirty sample sets which contain ten observations data of striation spacing for each sample, Fig. 13 shows the result of striation spacing analysis at 4000 µm from the crack initiation site in the two Al alloys. The mean of striation spacing means µx and standard deviation for the mean of striation spacing distribution σx are shown in the normal distribution graph in Fig. 13(a). The average striation spacing for Al6061-Zr alloy was calculated to be 1.07 µm, while the corresponding value for Al6061-Zr-Mg alloy was calculated to be 0.92 µm. It implies that the FCGR in Al6061-Zr alloy is faster than that in Al6061-Zr-Mg alloy. Fig. 13(b) shows the probability plots of two Al alloys which correspond to the normal distribution graphs in Fig. 13(a). The data of the means striation spacing in both Al alloys are approximately symmetrical. Besides, nearly sixty-eight percent data of the means striation spacing are distributed within one standard deviation of the mean. Thus, it can be concluded that three hundred data of striations spacing are enough for the present analysis to make the results consistent with the normal distribution. According to the central limit theorem, by increasing the number of observations, although the mean value of the sample means does not significantly change, the shape of the bell curve is narrower. Fig. 14 shows SEM photographs of two matching areas on both sides of the fracture surfaces of Al6061-Zr and Al6061-Zr-Mg alloys, indicated by boxes and arrows, respectively. Through careful observation, we found that striation formation in both Al alloys exhibited peak-to-valley matching, as shown in Fig. 14(a) and (b). The striation profile in Al6061-Zr alloy is more distinct than that in Al6061-Zr-Mg alloy. The striation areas in both Al alloys were also examined. For this purpose, the striation ratios were analyzed to determine the effect of excess Mg on generating the Mode I crack in the propagation region. On the basis of topography observations (Fig. 15), it was found that Al6061-Zr-Mg alloy generated the Mode I crack more easily and with a larger fatigue striation formation area compared to Al6061-Zr alloy. Fig. 16 shows an overview of the fracture surface in both Al alloys. In general, fracture surfaces can be classified into four types, namely striation, slant surface, smooth surface, and dimple, which are illustrated as 1, 2, 3 and 4, respectively, in Fig. 16(a). The development of stable crack growth is influenced by slant cracks. Strain localization is believed to play a significant role in determining the crack growth behavior of Al alloys. For Al6061-Zr-Mg alloy, in the presence of a high local stress in front of a crack tip, further damage that accumulates as a result of local plastic deformation is limited by the work hardening characteristics of the matrix. However, for a low work hardening material such as Al6061-Zr alloy, the presence of the damage zone under a sufficiently high local stress at the crack tip may lead to a localized work-softened zone, resulting in strain localization and shear or slant cracks [36]. Fig. 16(b) and (d) show the slant crack areas in Al6061-Zr alloy and Al6061-Zr-Mg alloy, respectively. In general, this phenomenon induces a local mixed- mode at the crack front, even if only a purely remote Mode I loading is applied, and it may lead to crack growth instability. Moreover, in the absence of striation, fractography observations indicate that the smooth surface is the dominant area on the fracture surface. It is believed that interaction of the crack front with particles or grain boundaries in the matrix structure cause the crack front to deflect from the straight growth direction, resulting in heterogeneity and smoothness of the crack surface. In some cases, the fracture surface is smooth and nearly flat, as shown in Fig. 16(c). In practice, flat fracture is dangerous because it is associated with a highvelocity crack growth [37]. Thus, stresses are maintained at high levels and fast propagation is promoted. As the locations of these types of fracture surfaces are random, they may induce scatter in the FCGR. Furthermore, dimples were observed in some areas on the fracture surface, indicating ductile fracture and rapid crack propagation. The dimple pattern was confirmed by observing both sides of the fracture surfaces that exhibited valley-to-valley matching. Fig. 16(e) shows the profile development of fatigue striation. The striations are

Fig. 7. S-N curves of the Al alloys.

Fig. 8. Fatigue crack length plotted against number of cycles. These curves indicate fatigue crack non-propagation in the Al6061-Zr-Mg alloy.

Fig. 9. Crack length as a function of number of cycles at a stress amplitude of 200 MPa.

deformation. Fig. 11 compares the fatigue striation spacing of the two Al alloys at 4500 µm from the crack initiation site. In addition, the mean value of striation spacing of each Al alloy is plotted as a function of the distance from the crack origin in Fig. 12. The striation spacing is closely related to the FCGR [32–34]. Examination of the fatigue striations indicated that the FCGR in Al6061-Zr-Mg alloy propagates slower than that in Al6061-Zr alloy. This finding is in good agreement with that of a 6

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Fig. 10. Fatigue crack growth behavior of the main cracks observed using plastic replica technique: (a) Al6061-Zr alloy and (b) Al6061-Zr-Mg alloy.

alloy were determined to be 19.5 and 111.2 nm, respectively. Thus, the fatigue striations in Al6061-Zr alloy were deeper than those in Al6061Zr-Mg alloy. On the basis of these findings, we propose different mechanisms of fatigue crack propagation in the two precipitation-hardened Al alloys with different Mg contents. These mechanisms are discussed in the next section.

characterized by ripple patterns [38], and the fracture surface is generally perpendicular to the loading direction. To investigate the profile development during the Mode I crack propagation, the formation of fatigue striation in Al6061-Zr-Mg alloy was examined by comparing its surface roughness with that of Al6061Zr alloy. Fig. 17 shows 3D SEM images of fatigue striation formation in the two Al alloys. From the surface roughness data shown in Fig. 18, it can be seen that the profile development of fatigue striation in Al6061Zr-Mg alloy is comparatively lower than that in Al6061-Zr alloy. The roughness measures Ra and Rz for Al6061-Zr-Mg alloy were determined to be 17.4 and 70 nm, while the corresponding values for Al6061-Zr

4. Discussion First, it should be noted that the addition of excess Mg produces a serrated flow, as shown in Fig. 6. The serrated flow is attributed to the 7

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Mean value of striation spacing, S (μm)

Fig. 11. Fracture surfaces with different striation spacing at 4500 µm from the crack initiation site: (a) Al6061-Zr alloy and (b) Al6061-Zr-Mg alloy.

Fig. 13. Analysis of the striation spacing at 4000 µm from crack initiation site: (a) Normal distribution, and (b) normal probability plot.

Fig. 12. Investigation of the striation spacing of the two Al alloys, which corresponds to the result in Fig. 9.

alloy, as shown in Fig. 8. The detailed mechanism of the fatigue crack non-propagation, which is related to the occurrence of DSA, has been discussed in previous studies [9,10]. In the present study, we focused on the microstructurally large fatigue cracks in the newly developed Al alloys rather than on the microstructurally small cracks, which were confirmed not to propagate at the fatigue limit. Through examination of the specimen surface, the crack in Al6061Zr-Mg alloy was observed to propagate more slowly than that in Al6061-Zr alloy. This trend becomes more distinct with increasing crack length, as shown in Figs. 9 and 10. The occurrence of stable Mode I crack growth can be considered after 2 × 105 cycles, in which the slower crack growth rate in Al6061-Zr-Mg alloy was supported locally by the smaller striation spacing compared to the case of Al6061-Zr alloy as shown in Fig. 11. However, it is important to note here that the result in Fig. 12 corresponds to the finding in Fig. 9, in which the result is referring to the mean value of the sample based on the limited number of the test specimen. Therefore, the analysis of striation spacing is

DSA caused by the excess Mg [39–41]. The serrated flow is presented due to the repeated dynamic pinning and unpinning effects. It was also reported that the serration can be occurred by the competition between DSA and shearing of precipitates by dislocations [42]. Moreover, in contrast to previous results obtained in the absence of Zr [9], the tensile strength of Al6061-Zr-Mg alloy is nearly equal to that of Al6061-Zr alloy. This is because of the suppression of grain-coarsening by the added Zr, as indicated by the light micrographs in Fig. 1. Accordingly, the fatigue limit of Al6061-Zr-Mg alloy is comparable to that of Al6061Zr alloy. Here, we note an important feature of the S-N relationship for Al6061-Zr alloy, i.e., the occurrence of failure after 107 cycles, as indicated in Fig. 7. By contrast, Al6061-Zr-Mg alloy does not fail, even after 5 × 107 cycles. This attribute of the latter alloy is afforded by the excess Mg and has been observed previously [9]. The difference between the fatigue lives at the fatigue limit of the two Al alloys stems from the occurrence of fatigue crack non-propagation in Al6061-Zr-Mg 8

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Ratio, AS / AO (%)

Fig. 14. Striation formation viewed from the other side of the fracture surface at 3500 μm from the crack origin: (a) Al6061-Zr alloy and (b) Al6061-Zr-Mg alloy.

supported by the findings of a previous study [43], in which the dominance of Mode I crack growth over Mode III crack growth was shown to cause the crack growth rate to decrease and the fatigue life of the specimen to increase. Moreover, as the fatigue crack propagates via a repeating process of plastic blunting and re-sharpening, the relationship between FCGR (dl/dN) and the crack tip opening displacement (ΔCTOD) is proportional [44], and for the plane strain condition ΔCTOD is defined as

AS : Area with striation in AO AO : Whole area

φ=

C (1 − v 2) K 2 Eσy

(3)

where C, K, v, E and σy are a constant, the stress intensity factor, Poisson’s ratio, Young’s modulus, and the yield strength, respectively. According to Dugdale’s model [45], for a remote stress σ∞ acting on an infinite body with a crack size of 2c, the plastic zone size ωp is expressed as

Fig. 15. Ratio of area with striation to whole specified area.

πσ ⎞ ⎛ ωp = c ⎜sec ⎛⎜ ∞ ⎞⎟ − 1⎟. 2σ ⎠ ⎝ ⎝ y⎠

expanded using the central limit theorem to determine an approximate value of the population mean, in which the result can be more useful in terms of material properties. There are two points noteworthy to be considered from the result shown in Fig. 13. Firstly, the average striation spacing in Al6061-Zr-Mg alloy was lower than that in Al6061-Zr alloy, which indicates that the slower FCGR in Al6061-Zr-Mg alloy compared to Al6061-Zr alloy. Secondly, it is shown that the distribution of striation spacing data in Al6061-Zr alloy disperses more than those in Al6061-Zr-Mg alloy. It implies that there was wider variation in the striation spacing data in the Al6061-Zr alloy than in the Al6061-Zr-Mg alloy, which might be attributed to the unstable crack growth behavior in Al6061-Zr alloy compared to Al6061-Zr-Mg alloy because of less tendency in generating Mode I crack. The retardation of the FCGR in Al6061-Zr-Mg alloy might be due to the high tendency of striation formation, as indicated in Fig. 15. The authors believe that the Mginduced work hardening will generate a large number of active slip systems, resulting in more non-localized slip and a large area of striation formation. Consequently, this phenomenon probably induces a highly stable Mode I crack growth in Al6061-Zr-Mg alloy and possibly facilitates crack closure, such as the asperity-induced type. This result is

(4)

The new aluminum alloy with excess Mg has a higher work hardening rate, which results in a smaller plastic zone size compared to Al6061-Zr alloy. Eq. (4) gives the inverse correlation between ωp and σy . According to Eq. (3), ΔCTOD depends on σy , which influences the FCGR of microstructurally large cracks. Hence, it is reasonable that Al6061Zr-Mg alloy has a slower FCGR compared to Al6061-Zr alloy. On the basis of the above-mentioned findings, the effect of DSA on the microstructurally large cracks was examined locally by investigating the fatigue striation morphology. DSA is induced by the interaction between solute Mg and the matrix dislocations in Al6061Zr-Mg alloy, resulting in greater resistance to plastic deformation. In other words, the mobility of the dislocations during crack growth is restricted by the DSA effect. This condition easily generates a pair of alternating slips with many slip traces in Al6061-Zr-Mg alloy. Furthermore, the high deformation resistance of Al6061-Zr-Mg alloy results in a smaller ΔCTOD during the blunting process, and hence a smaller striation spacing compared to Al6061-Zr alloy (an example is shown in Fig. 11). In addition, Al6061-Zr alloy has a deeper striation 9

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Fig. 16. General view of fracture surfaces: (a) schematic classification of fracture surfaces; (b) fracture surface of Al6061-Zr alloy at 3500 µm from the crack initiation site; (c) profile development of the fracture surface, which corresponds to the scanning line in Fig. 15(b); (d) fracture surface of Al6061-Zr-Mg alloy at 3500 µm from the crack initiation site; and (e) profile development of the fracture surface, which corresponds to the scanning line in Fig. 15(d).

work hardening of the above-mentioned Al alloys: (i) dislocation-dislocation interaction, (ii) dislocation-precipitate interaction, and (iii) dislocation-(solute Mg) interaction. It is considered that mechanism (i) significantly affects Al-Mg alloy, in which the excess Mg content increases the dislocation density [46]. Hence, Al6061-Zr-Mg alloy has a higher dislocation density compared to Al6061-Zr alloy, and the number of dislocations further increases with deformation. Consequently, the distance between the dislocations decreases, resulting in restricted dislocation motion. In other words, the plastic resistance of Al6061-Zr-Mg alloy increases with increasing the dislocation density, which contributes to the rapid strain hardening phenomenon. With regard to mechanism (ii), the work hardening can also arise from interactions between the mobile dislocations and the precipitates. However, we assume that all the precipitates can be cut by the dislocations during severe plastic deformation near the crack tip in both Al alloys. Nevertheless, the dislocation-precipitate interactions can be influenced by mechanism (iii). It is known that work hardening is also induced by interactions between the glide dislocations and the solute Mg [47]. The distinct fatigue striation in Al6061-Zr alloy is attributed to the high tendency of local plastic deformation due to the reduced resistance to dislocation motion. It is considered that the absence of solute Mg in the matrix exposes the alloy to slip localization, which is associated with easy shearing of precipitates by dislocation motion [48,49]. When a dislocation cuts through a precipitate, the movement of subsequent dislocations on the same plane becomes easier and local softening is promoted. This phenomenon is in good agreement with previous studies on samples taken from Al alloys [50] that exhibit greater material instability when the material is subjected to a complex load.

Fig. 17. 3D images of different fatigue striation developments at 4000 μm from the crack origin: (a) Al6061-Zr alloy and (b) Al6061-Zr-Mg alloy.

profile compared to Al6061-Zr-Mg alloy, which is believed to be due to the greater work hardening of the latter. Three reasonable mechanisms are thought to generally affect the 10

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Fig. 18. Surface roughness analysis: (a) 2D SEM image of Al6061-Zr alloy, (b) 2D SEM image of Al6061-Zr-Mg alloy, and (c) graphical roughness profiling.

dislocations, as shown in Fig. 19(a)(iii) and (iv). This condition is due to the low resistance to dislocation motion, and it may also be promoted by the high tendency of precipitate shearing and local softening [54], leading to highly localized deformations, as shown in Fig. 18(c). By contrast, DSA induces plastic resistance and greater work hardening through interactions between the gliding dislocations on the slip system and the solute Mg. When the tensile load is increased, the dislocations begin to move on the first slip plane, as shown in Fig. 19(b)(iii). As the shear resistance of this plane increases, the slip activity continues on the second slip plane in response to the resolved shear stress, as shown in Fig. 19(b)(iv). The dislocations on this plane are blocked in the same manner as that in the case of the first slip plane owing to the greater work hardening in this alloy. This facilitates the occurrence of multiple pairs of alternating slips, with limited numbers of dislocations gliding along the slip planes owing to the DSA effect (see Fig. 19(b)(v) and (vi)). The numbers of dislocations do not differ significantly among the planes as can be observed from Fig. 19(b)(iii–vi). This phenomenon leads to the blunting of cracks that are terminated early, and ΔCTOD becomes smaller compared to that in the absence of the DSA effect. When the maximum load is reached, the load is reversed, and the positive dislocations begin to move closer to the crack surface. We have assumed that the positive dislocations cannot disappear through the crack surface owing to the presence of a thin oxide layer formed by exposure of the new surface to the environment [55,56]. This promotes slip irreversibility, which resists the returning dislocations. It is noteworthy that the blunting of the crack tip also causes slip irreversibility [57]. Consequently, during the unloading process, dislocation pairs nucleate from Frank-Read sources in the parallel slip planes. A dislocation pair consists of two types of dislocations with their Burgers vectors in opposite directions. In both Al alloys, the positive dislocations move towards the crack surface and are eliminated as they approach the surface, while the negative dislocations persist (see Fig. 19(a)(v–viii) and (b)(vii–x)). However, the crack closure behaviors in the two Al alloys during unloading are different. The alloy without DSA has a higher ratio of the average striation height to the striation spacing. This implies that the reverse slip in this alloy during unloading occurs in a

Furthermore, the occurrence of DSA in Al6061-Zr-Mg alloy causes the solute Mg to act as an obstacle to dislocation motion, with the solute Mg interacting with the mobile dislocations and locally pinning them. Consequently, easy movement of the dislocations is hindered, and this constitutes a barrier to subsequent dislocations. Hence, the frequency of precipitates being cut by dislocations is reduced, and the action of pinning dislocations presumably increases the work hardening. It is noted that the presence of extra solute Mg in the matrix of Al6061-ZrMg alloy prevents susceptibility to slip localization on particular slip planes. The proposed effect of DSA on the fatigue crack growth model is based on the observation of the fatigue crack growth behavior on the specimen surface and fractographic analysis of the fracture surface. The two considered Al alloys exhibited similar patterns of striation formation, with peak-to-valley matching, as shown in Fig. 14. This means that alternating slip systems are activated during crack growth. Similar patterns of fatigue striation formation were reported in [14,51]. Another striation morphology is peak-to-peak or valley-to-valley matching [31,52,53]. Careful observation of the fracture surface revealed no change in the general striation formation over the next few cycles. Thus, the peak-to-valley cycle of striation formation is more reasonable because the peak-to-peak cycle increases the additional residual stress owing to the reduction of the material volume. This would result in an increase in the compressive residual stress with increasing the number of cycles and hence continuous reduction of the applied stress and prevention of fatigue crack propagation. Fig. 19(a) and (b) show the fatigue crack growth models of the Al alloys without and with DSA, respectively. The letters and numbers in the dislocation models indicate the positions in the load cycle, as shown in Fig. 19(a)(i) and (b)(i). In general, the slip is alternatively activated in a concentrated manner during loading and distributed during unloading. During the loading process, dislocations are emitted from the crack tip owing to stress concentration. The phenomenon occurs in both Al alloys. However, in the absence of DSA, the crack opens with a pair of concentrated slips on slip planes a and b. A large number of dislocations emanate from the crack tip and glide along the slip planes, with a pair of concentrated slips alternately activated by the avalanche [14] of

11

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Fig. 19. Dislocation model of Mode I fatigue crack growth mechanism in precipitation-hardened Al alloys: (a) Al6061-Zr alloy and (b) Al6061-Zr-Mg alloy.

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(6) As determined by fractographic analysis, the fatigue striation formation area in Al6061-Zr-Mg alloy is larger than that in Al6061-Zr alloy, which may be credited with the stable Mode I crack growth behavior in Al6061-Zr-Mg alloy.

concentrated manner rather than a distributed manner, with a limited number of slip planes activated during one cycle, as shown in Fig. 19(a) (v–viii). Gliding dislocations initially occur on the plane adjacent to the slip plane that was activated last during loading, identified as slip plane 1 in Fig. 19(a)(v), and this is followed by slip plane 2. Subsequent highly localized deformations occur progressively on slip planes 3, 4, 5, 6, 7, 8, 9 and 10 owing to the low resistance to dislocation slip (see Fig. 19(a)(vi–viii)). This situation constitutes a distinct type of striation formation with a high roughness profile, as shown in Fig. 18(c). By contrast, the presence of DSA in the alloy induces the activation of many slip planes, with slip steps occurring during loading. The reverse slip during unloading is thus distributed, beginning on slip planes 1 and 2, as shown in Fig. 19(b)(vii). Then, the crack is sequentially closed by the alternating slip occurring on each slip plane, as shown in Fig. 19(b) (viii–x). This condition promotes the formation of shallow fatigue striations with a low roughness profile, as shown in Fig. 18(c), resulting in highly stable crack growth in the Al alloy with DSA. In a future study, we will investigate whether the FCGR scatter of microstructurally large cracks might be because of the stability of the Mode I crack growth; stability in this context refers to the ability to generate a large area of fatigue striation formation. Finally, it should be noted that the experimental work in this study was conducted at room temperature. It is known that the DSA effect can be influenced by the temperature [58] and strain rate [59]. In general, DSA occurs slowly at room temperature and is accelerated at higher temperatures, because diffusion of the solute atoms is enhanced by increasing the temperature. It is expected that at higher temperatures, the effect of DSA in Al6061-Zr-Mg alloy becomes more significant, where the pinning of dislocations by solute Mg occurs rapidly during straining, and this phenomenon increases the stress required for the dislocation movement. Consequently, work hardening occurs easily on the primary slip plane and the secondary slip is activated instead. The rapid process of work hardening on the slip plane results in shallower fatigue striation formation, increasing the striation ratio; thus, crack growth in Al6061-Zr-Mg alloy is highly stable. Moreover, it is believed that the plastic zone evolution will be suppressed by increasing the work hardening in the Al alloy with DSA, resulting in enhanced fatigue resistance [60].

It is speculated that the stable Mode I crack growth in Al6061-Mg alloy induces smaller scatter of the FCGR of microstructurally large cracks. Further investigation is required to understand the effects of DSA on the FCGR scatter of the two considered Al6061-T6 alloys. This issue will be addressed elsewhere. A deeper understanding of the crack growth behaviors of Al alloys containing excess Mg would facilitate enhancement of their fatigue properties. Acknowledgement One of the authors (S.F. Anis) gratefully acknowledges the financial support provided by the Government of Malaysia and the University of Technology, Malaysia. This work was financially supported by JSPS KAKENHI (JP 16H06365). References [1] G.S. Cole, A.M. Sherman, Lightweight materials for automotive applications, Mater. Charact. 35 (1995) 3–9. [2] W.S. Miller, L. Zhuang, J. Bottema, A.J. Wittebrood, P. De Smet, A. Haszler, A. Vieregge, Recent development in aluminum alloys for the automotive industry, Mater. Sci. Eng., A 280 (2000) 37–49. [3] J.R. Davis, Aluminum and aluminum alloys, ASM Specialty Handbook, ASM International, 1993. [4] B.F. Jogi, P.K. Brahmankar, V.S. Nanda, R.C. Prasad, Some studies on fatigue crack growth rate of aluminum alloy 6061, J. Mater. Process. Technol. 201 (2008) 380–384. [5] R.C. Rice, J.L. Jackson, J. Bakuckas, S. Thompson. Metallic materials properties development and standardization (MMPDS) Handbook, NTIS Virginia Scientific Report. 2003. [6] Aluminum Company of America, Alcoa Structural Handbook, Alcoa, Pittsburgh. 1960. [7] M. Goto, N. Kawagoishi, H. Nisitani, A. Miura, Statistical investigation of small crack growth rate in age-hardened Al alloy 6061–T6, Trans. JSME. 62–595 (1996) 671–677. [8] S. Kalpakjian, Manufacturing Engineering and Technology, 3rd ed., Addison-Wesley Publishing Co., 1995. [9] T. Shikama, Y. Takahashi, L. Zeng, S. Yoshihara, T. Aiura, K. Higashida, H. Noguchi, Distinct fatigue crack propagation limit of new precipitation-hardened aluminum alloy, Scr. Mater. 67 (2012) 49–52. [10] L. Zeng, T. Shikama, Y. Takahashi, S. Yoshihara, T. Aiura, H. Noguchi, Fatigue limit of new precipitation-hardened aluminum alloy with distinct fatigue crack propagation limit, Int. J. Fatigue 44 (2012) 32–40. [11] S.F. Anis, M. Koyama, H. Noguchi, Investigation on mode I propagation behavior of fatigue crack in precipitation-hardened aluminum alloy with different Mg content, Mater. Sci. Forum 889 (2017) 143–147. [12] R.O. Ritchie, J. Lankford, Small fatigue cracks: A statement of the problem and potential solutions, Mater. Sci. Eng.. 84 (1986) 11–16. [13] K. Tokaji, T. Ogawa, Y. Harada, Z. Ando, Limitations of linear elastic fracture mechanics in respect of small fatigue cracks and microstructure, Fatigue Fracture Eng. Mater. Struct. 9 (1986) 1–14. [14] P. Neumann, Coarse slip model of fatigue, Acta Metall. 17 (1969) 1219–1225. [15] E. Pink, A. Grinberg, Stress drops in serrated flow curves of Al5Mg, Acta Metall. 30 (1982) 2153–2160. [16] R.B. Schwarz, L.L. Funk, Kinetics of the Portevin-Le Chatelier effect in Al 6061 alloy, Acta Metall. 33 (1985) 295–307. [17] R. Kral, P. Lukac, Modelling of strain hardening and its relation to the onset of Portevin- Le Chatelier effect in Al-Mg alloys, Mater. Sci. Eng., A 234 (1997) 786–789. [18] W. Wen, Y. Zhao, J.G. Morris, The effect of Mg precipitation on the mechanical properties of 5xxx aluminum alloys, Mater. Sci. Eng., A 392 (2005) 136–144. [19] M.N. Rittner, J.R. Weertman, J.A. Eastman, K.B. Yoder, D.S. Stone, Mechanical behavior of nanocrystalline aluminum-zirconium, Mater. Sci. Eng., A 237 (1997) 185–190. [20] W. Yuan, Z. Liang, Effect of Zr addition on properties of Al-Mg-Si aluminum alloy used for all aluminum alloy conductor, Mater. Des. 32 (2011) 4195–4200. [21] F. Wang, D. Qiu, Z.-L. Liu, J.A. Taylor, M.A. Easton, M.-X. Zhang, The grain refinement mechanism of cast aluminum by zirconium, Acta Mater. 61 (2013) 5636–5645. [22] H. Li, Z. Li, Y. Liu, H. Jiang, Effect of zirconium on the microstructure and mechanical properties of Zn-4%Al hypoeutectic alloy, J. Alloy. Compd. 592 (2014) 127–134. [23] G.C. Sih, Multiscale Fatigue Crack Initiation and Propagation of Engineering Materials: Structural Integrity and Microstructural Worthiness, Springer, 2008.

5. Conclusions In this study, the fatigue crack growth characteristics of microstructurally large fatigue cracks in newly developed Al6061-Zr and Al6061-Zr-Mg alloys were experimentally examined using smooth specimens of the alloys. The fatigue crack propagation mechanisms of the alloys were observed to be affected by their excess Mg contents compared to other Al6061-T6-based alloys. The main conclusions that can be drawn from the experimental results are as follows: (1) Al6061-Zr-Mg alloy exhibits a distinct fatigue limit, whereas Al6061-Zr alloy does not. (2) The striation spacing in Al6061-Zr-Mg alloy is smaller than that in Al6061-Zr alloy; hence, the FCGR in the former is lower than that in the latter. Crack growth retardation in Al6061-Zr-Mg alloy is speculated to be due to its easier in generating Mode I cracks. (3) Fatigue striation examination confirmed the enhancement of Mode I crack growth in Al6061-Zr-Mg alloy. The higher tendency of this alloy to generate Mode I cracks may be attributed to the occurrence of DSA. (4) The profile development of fatigue striations in the two investigated Al alloys exhibits peak-to-valley matching rather than symmetrical matching. (5) Al6061-Zr-Mg alloy has a shallower fatigue striation profile compared to Al6061-Zr alloy. This is believed to be due to the constrained dislocation gliding during crack growth, attributed to the DSA effect. 13

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