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Microstructures and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy with addition of rare earth Er
Materials Science and Engineering A 527 (2010) 1806–1814
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Microstructures and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy with addition of rare earth Er Song Bai a,b , Zhiyi Liu a,b,∗ , Yuntao Li a,b , Yanhui Hou a,b,c , Xu Chen a,b a b c
Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Central South University, Changsha 410083, People’s Republic of China School of Material Science and Engineering, Central South University, Changsha 410083, People’s Republic of China School of Mechanics and Engineering, Southwest Jiaotong University, Chengdu 610031, People’s Republic of China
a r t i c l e
i n f o
Article history: Received 17 March 2009 Received in revised form 31 October 2009 Accepted 4 November 2009
Keywords: Al–Cu–Mg–Ag alloy Microstructures Fatigue fracture behavior Er
a b s t r a c t The effects of rare earth erbium (Er) on microstructures and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy were investigated. Microstructural examinations first revealed that the precipitation kinetics of phase was distinctly retarded by promoting the formation of phase with the Er addition during the initial aging. The fatigue crack propagation resistance of Er-containing microstructure was significantly enhanced arising from the presence of the crystallographic secondary cracks, which was directly relative to the large grain size. Results also suggested that the dendritic substructure of as-cast Al–Cu–Mg–Ag alloy was refined remarkably by Er addition. © 2009 Elsevier B.V. All rights reserved.
1. Introduction The addition of Ag with small amount is well known to enhanced the aging hardening response in a wide range of Al–Cu–Mg alloys by the precipitation of the fine and uniform plates termed phase on the {1 1 1}␣ planes [1,2]. phase is believed to be a variant of equilibrium phase (Al2 Cu) and keeps stable at elevated temperature arising from the high coarsening resistance [3]. In addition to the attractive strength properties [4], compared with commercial aluminium alloys, higher creep resistance was also exhibited if the alloys are in the underaged conditions [5–8]. These above make Al–Cu–Mg–Ag-based alloys the promising materials for aerospace applications. With the development of the aerospace industries, mechanical properties of Al-based alloys are required to be further improved to satisfy the applications that used in tension-dominated aircrafts require excellent fatigue damage tolerance. Furthermore, it is well accepted that small additions of rare earth elements are of great importance to improve the mechanical properties of Al-based alloys and provide extended alloying capabilities. Over the past several years, in the case of Al-based alloys rare earth Scandium (Sc) has been most commonly studied as beneficial microalloying element
∗ Corresponding author at: School of Material Science and Engineering, Central South University, Lu Mountain South Road, Hunan, Changsha 410083, People’s Republic of China. Tel.: +86 731 88836011; fax: +86 731 88876692. E-mail address: [email protected] (Z. Liu). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.11.011
[9–11], suggesting the beneficial effects on mechanical properties that associated with remarkably grain refinement, dynamic recrystallization controlling and also fatigue as proposed by Wirtz et al. [12]. However, the applications of Sc-containing alloys are extremely restricted due to the high cost of Sc addition. Instead, much attention was focused on rare earth Erbium (Er), which is much cheaper than Sc. A previous work conducted by Xu et al. [13] revealed that, with the addition of Er, the mechanical properties of Al–Mg alloys were strongly improved by the presence of the fine dispersed Al3 Er particles after heat treatment, whereas no similar effects were observed in a hot-extruded Al–Mg alloy without solution treatment as suggested by Wu et al. [14]. Based on the research by Karnesky et al. [15], the volume fraction of precipitates in the Al–0.08 at.% Sc alloy increased with small addition of 0.02 at.% Er. In Al–Zn–Mg alloys, Xu et al. [16] demonstrated that Er improves the strength considerably by precipitation strengthening mechanisms and grain refinement and retards recrystallization due to the pinning effect of fine dispersed A13 Er precipitates on dislocations and subgrain boundaries. Li et al. [17] investigated the influences of Er on the grain refinement in an as-cast Al–Cu–Mg–Ag alloy and found that Er atoms may segregate at grain boundary during solidification and resulted in ternary Al8 Cu4 Er phase. Besides, rare earth Cerium (Ce) and Ytterbium (Yb) also attract much attention. The addition of Ce promotes the precipitation and improves the thermal stability of the phase in Al–Cu–Mg–Ag alloy [18,19]. Similar results were also observed with addition of Yb [20]. Studies of the fatigue behavior of the Al–Cu–Mg–Ag-based alloys are limited [21], especially in the case of microalloying with rare
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Table 2 Room temperature mechanical properties of underaged Al–Cu–Mg–Ag–(Er) alloys.
Table 1 Compositions of the studied Al–Cu–Mg–Ag–(Er) alloys (in wt.%). Cu Er-free Er-containing
6.45 6.21
Mg 0.65 0.62
Ag 0.46 0.44
Mn 0.29 0.28
Ti 0.09 0.09
Zr 0.14 0.15
Er – 0.32
Fe 0.04 0.05
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Si 0.01 0.01
Al Bal. Bal.
earth Er. The present work mainly aims at clarifying the influence of rare earth element Er on the precipitation evolution and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy during the initial aging. The mechanism governed by Er was also proposed from the obtained results.
Alloys
Temperature ◦
b (MPa)
0.2 (MPa)
ı (%)
Er-free
165 C × 2 h 165 ◦ C × 4 h
486.1 500.6
422.6 476.4
16.2 11.3
Er-containing
165 ◦ C × 2 h 165 ◦ C × 6 h
453.7 488.2
341.9 460.1
21.0 11.7
condition), the FCP data was considered as a reference for further design. 2.3. Microstructural analysis
2. Experimental procedures 2.1. Sample preparation and heat treatment Al–Cu–Mg–Ag alloys with high Cu/Mg ratios are used in this study and have the compositions as listed in Table 1. It should be noted that although the relative amount of Cu in Al–Cu–Mg–Ag–Er alloy is a little lower as compared to Al–Cu–Mg–Ag alloy, it has no significant influence on the experimental results. Both alloys were cast into a cast iron mould. The ingots of Al–Cu–Mg–Ag–(Er) alloys experienced homogenization followed by air cooling to room temperature, scalped and hot roll to 2.2 mm thick strips at about 450 ◦ C. The alloy samples were solution treated at 515 ◦ C for 6 h, cold water quenched and immediately aged at 165 ◦ C in air furnace for various times. 2.2. Tensile and FCP testing Specimens for tensile testing were prepared vertical to the longitudinal direction of the strips with gauge length 30 mm and 2.0 mm thick. Tensile testing was performed at room temperature on the CSS-44100 type testing machine with 2 mm/min loading speed. Fatigue crack propagation (FCP) testing was performed on compact tension (CT) specimens taken from the strips in the L–T orientation with a size (in mm) of 47.5 × 45 × 2 (L × W × B) to obtain the fatigue crack growth (FCG) rates that are above 10−5 mm/cycle. The underaged samples were chosen for FCP testing. All FCP tests were conducted at a stress ratio (R = Kmin /Kmax ) of 0.1 with a loading frequency of 10 Hz on an Instron 8801 50 kV fatigue tester at room temperature, laboratory air environment. The crack length was monitored by elastic compliance technique. Though the plane stress condition was inevitably present during FCP testing and not in accordance with the standard of fatigue test, here it is necessary to pointed out that the purpose of this study was to clarify the fatigue behavior of Er-containing alloy in thin sheet which is closest to its condition in real applications (that is plane stress
Optical metallographic examination was used to reveal the grain structures of both alloys. Transmission electron microscopy (TEM), along with selected area electron diffraction (SAED), was employed to characterize the differences of the microstructure evolution in Al–Cu–Mg–Ag alloys with and without Er during aging treatment. Samples suitable for TEM were thin slices with 3 mm in diameter and electro-polished by using twin-jet equipment with a voltage of 10 V in a 70% methanol and 30% nitric acid at approximately −20 ◦ C, after which the slices were cleaned in ethanol at room temperature for at least 5 min. These were then examined on a Tecnai G2 20 ST TEM machine operating at 200 kV. Fatigue fracture surfaces of the samples cyclically deformed in FCP testing was analyzed by KYKY 2800 scanning electron microscopy (SEM) and the distribution of rare earth Er was measured by Quanta 200 machine. Three-dimensional atom probe (3DAP) technique was also used to reveal how the addition of Er influences the precipitation reactions in Al–Cu–Mg–Ag alloy. 3. Results 3.1. Microstructures and mechanical properties The typical grain structures of the as-cast alloys used in this study were characterized by optical microscopy in order to reveal the influence of Er addition as shown in Fig. 1. Fig. 1(a) illustrates the cast structure of Al–Cu–Mg–Ag alloy that predominated by large equiaxed grains with coarsening dendritic substructure. The grain boundaries were decorated by continuous phases. Great branch distance was present and the average grain size of this casting was about 0.4 mm. Apparently, the distance between the branches of the dendritic substructure was greatly reduced as a direct consequence of the addition of Er shown in Fig. 1(b), which is consistent with earlier observations [17] on similar materials. The grain was slightly refined and was about 0.2 mm. Grain structures of the underaged samples used in the FCP testing were shown in Fig. 2,
Fig. 1. Typical grain structures of as-cast alloys (a) Al–Cu–Mg–Ag and (b) Al–Cu–Mg–Ag–Er.
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Fig. 2. Typical grain structures of the FCP testing samples (a) Al–Cu–Mg–Ag alloy aged at 165 ◦ C for 4 h and (b) Al–Cu–Mg–Ag–Er alloy aged at 165 ◦ C for 6 h.
which both show the fully recrystallised grains. The Er-containing alloy exhibits a much larger recrystallised grain structure. The microstructures of underaged Al–Cu–Mg–Ag–(Er) alloys were characterized as shown in Fig. 3, corresponding room temperature mechanical properties are given in Table 2. The microstructure of Al–Cu–Mg–Ag alloy dominated by phase formed on the {1 1 1}␣ planes, together with fine plates on the {0 0 1}␣ planes, was found after aging for 2 h at 165 ◦ C as shown in Fig. 3(a), which was consistent with SAED pattern. Aging for 4 h at 165 ◦ C mostly produced phase and phase was only present as a minor phase as seen in Fig. 3(c), indicating the precipitation of phase was enhanced as aging goes at the expense of phase supported by previous study [22]. Corresponding SAED pattern in Fig. 3(c) was hard to reveal the spots that associated with phase, only reflections at the 1/3 and 2/3 {0 2 2}␣ positions produced by phase were observed. However, as shown in Fig. 3(b), with the addition of Er differences were distinctly observed in the appearance of phase compared with Fig. 3(a). Only phase was present and close examination of the microstructure in the thin TEM samples taken near 1 1 0␣ zone axis failed to reveal any precipitates that could be identified as phase. SAED pattern taken near 0 0 1␣ zone axis in Fig. 3(b) only confirmed the existence of phase. Considering the competition between and phases, it is suggested that the role of rare earth Er is to suppress the precipitation of phase by promoting the formation of phase at early stage of aging. The precipitation of phase was considerably accelerated following aged at 165 ◦ C for 6 h as revealed by the fine and uniform dispersions of phase along with plenty of fine precipitates observed in Fig. 3(d) and corresponding SAED pattern shown obvious reflections of and precipitates, although the intensity of phase was extremely weak, notably that the redissolving of phase occurred in Al–Cu–Mg–Ag alloy (seen in Fig. 3(c)) have been remarkably impeded. It is generally accepted that, in Al–Cu–Mg–Ag alloy, the precipitate kinetics of phase precipitates is greatly accelerated by Ag addition, since Ag and Mg atoms are strongly segregated to the broad face of phase as aging goes, lying on the {1 1 1}␣ planes, which reduces the misfit strain at the ␣/ interface and promotes the growth of phase [2,22,23]. Consequently, 3DAP was employed to clarify the microalloying behavior of Er addition. Despite the limitation of the field range, however, 3DAP result failed to reveal any evidence of Er atoms segregate to the ␣/ interface as seen in Fig. 4. This may attribute to the trace element of Er in this study.
Although Er atoms in as-cast Al–Cu–Mg–Ag alloy mainly segregate at grain boundary during solidification and resulted in ternary Al8 Cu4 Er phase [17], part of which are soluble in Al matrix during heat treatment, if not, the retardant of precipitation of phase is hard to explain because the Al8 Cu4 Er phase has no direct effect on the precipitation in Al matrix. Fig. 5 illustrates the backscattered electron (BSE) image of underaged Al–Cu–Mg–Ag–Er alloy (165 ◦ C × 6 h) and elements mapping of Cu and Er. It was clear that coarse Al8 Cu4 Er phase in as-cast structure was not visible and the distribution of Cu and Er was random. No obvious co-segregation of Cu and Er was present. This implies that Er atoms soluted in the Al matrix are associated with the slow precipitation response at the initial aging. 3.2. Fatigue fracture behavior As seen in Table 2 two studied underaged conditions (165 ◦ C × 4 h of Al–Cu–Mg–Ag alloy and 165 ◦ C × 6 h of Al–Cu–Mg–Ag–Er alloy) show very similar values of room temperature mechanical properties. Therefore, it is expected that different fatigue crack growth rates of two alloys are not resulted, at least, from the differences of mechanical properties. The variation of FCG rates with the stress intensity factor range (K) is given in Fig. 6 for the comparison between both underaged alloys. Apparently, both da/dN–K curves show the three typical stages from Kth up to the critical stress intensity amplitude where the final fatigue fracture occurs. Contrary to the expected above, even though no pronounced differences in FCG rates exist in nearthreshold regime (below K values of about 13 MPa m0.5 ), some clear deviation between the two curves was observed in the intermediate K region, Paris regime, where the Al–Cu–Mg–Ag alloy showed high FCG rate whereas higher fatigue crack propagation resistance against slow FCG rate was revealed for the microstructure of Al–Cu–Mg–Ag–Er alloy. It also should be emphasized that even at high K region (above K values of about 28 MPa m0.5 ), comparing with Al–Cu–Mg–Ag alloy which was obviously seen that FCG rate accelerated dramatically when K equals to 28 MPa m0.5 and then followed by fast fatigue fracture, the Er-containing alloy still exhibited higher fatigue crack propagation resistance and was insensitive to the variety of K, i.e., the FCG rate of Al–Cu–Mg–Ag alloy was several time higher than that of Er-containing alloy at K of 30 MPa m0.5 . The fatigue fracture of underaged Al–Cu–Mg–Ag–Er alloy occurred ultimately at about K of 42 MPa m0.5 . As revealed in Fig. 6, the crack propagation resistance was greatly enhanced with
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Fig. 3. TEM bright field micrographs and corresponding SAED patterns of the underaged Al–Cu–Mg–Ag alloy aged for (a) 2 h and (c) 4 h and Al–Cu–Mg–Ag–Er alloy aged for (b) 2 h and (d) 6 h at 165 ◦ C. Electron beam is approximately parallel to 1 1 0␣ zone axis in (a), (c) and (d) versus 0 0 1␣ zone axis in (b).
the addition of Er, especially in the high K values. It is assumed that the role of Er addition on the fatigue fracture behavior of the Al–Cu–Mg–Ag alloy should be concluded to stabilize the fatigue crack propagation as K increases. Examination of the fatigue fracture surfaces of both underaged alloys were carried out for comparisons at (i) low magnification to clarify the overall fatigue fracture morphology of different regimes and (ii) high magnification to identify the characteristics of the fatigue cracks growth and the fine-scale fatigue fracture features for both alloys in FCP testing. Representative fatigue fracture features were shown in Figs. 7–9 taken from fatigue fracture regions where the fatigue crack growth rates were measured. The near-threshold fatigue fracture surfaces of underaged alloys given in Fig. 7 shows different morphology. Compared with Fig. 7(a) and (b), the surface of Al–Cu–Mg–Ag alloy was characterized by minor fibrous beach marks that are roughly parallel to the prin-
cipal fatigue crack growth direction versus many main shearing ridges with the addition of Er were observed. According to the high magnification view shown in Fig. 7(c) and (d), the presence of microscopic voids arising from the localized micro-plastic deformation at the incoherent Mn/Zr-rich dispersoid particles through the matrix in Al–Cu–Mg–Ag alloy were found, which were not present in Er-containing alloy. It is generally accepted that the voids can be considered as being equivalent to second-phase particles having zero stiffness and facilitate a gradual increase of both the size and relative volume fraction of voids in cyclic deformation that ultimately result in the increments of crack-tip extension [24]. However, with the addition of Er, the fracture surface seems to be comprised of crystallographic shearing ridges joined by each other as revealed in Fig. 7(d). Crack deviation occurs apparently because the crack follows these weakened planes and deviates 20–40◦ with respect to the principal crack growth direction. FCG in this region
Fig. 4. 3DAP elemental mapping of a selected region of the Al–Cu–Mg–Ag–Er alloy aged for 6 h at 165 ◦ C.
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Fig. 5. BSE image and elements mapping of Al–Cu–Mg–Ag–Er alloy aged at 165 ◦ C for 6 h.
occurs by a transgranular shear mechanism. In this study, at low K values although the fatigue fracture surfaces did shown some differences as seen in Fig. 7, the FCG rates were almost the same for both alloys shown in Fig. 6. As shown in Fig. 8, it was observed that significant differences occurred in Paris regime between underaged Al–Cu–Mg–Ag and Al–Cu–Mg–Ag–Er alloys for which the most distinct deviation in FCG rates were measured as revealed in Fig. 6. The overall morphology of the fatigue fracture surfaces in Paris regime for both alloys was shown in Fig. 8(a) and (b), both exhibiting fairly well-developed fatigue characteristics and corresponding high magnification observations were present in Fig. 8(c)–(f), respectively. In Al–Cu–Mg–Ag alloy, the region of stable fatigue crack propagation comprised of pockets of fibrous shell marks formed during cyclic deformation and widely dispersed microscopic voids formed around dispersoid particles were observed. In contrast, few distinct microscopic voids were shown with the addition of Er. Fracture was predominantly transgranular with a random distribution of slip steps on which many regular coplanar cracks, termed secondary cracks, formed. Evidence of some ductile tearing arrowed in Fig. 8(b) was also exhibited. It can be clearly seen, by comparison between Fig. 8(c) and (d), that the transgranular fracture regions of Al–Cu–Mg–Ag alloy was predominantly comprised of a dispersion of microscopic voids and spherical shaped Mn/Zr-rich dispersoid particles of varying sizes, although some microscopic cracks were present as arrowed and circled in Fig. 8(c), whereas the presence of secondary cracks in a much higher density was visible in Er-containing alloy. Furthermore, the orientation of secondary cracks alignment was not along the principal fatigue crack growth direction but normal to the fracture surface. Fig. 8(e) and (f) illustrated some details of the fatigue fracture regions circled in Fig. 8(c) and (d), respectively. Secondary cracks were practically absent in Al–Cu–Mg–Ag alloy, only microscopic cracks were observed as arrowed in Fig. 8(e). Fig. 8(f) exhibits the presence of arrays of secondary cracks in Er-containing alloy and it was also obvious that minor uniformly spaced fatigue striations appear as arrowed in the surface of secondary cracks, which directly indicate the propagation process of secondary cracks also occurs during the cyclic deformation. SEM observations analysis revealed that the occurrence of secondary cracks results in a more complex crack path, thereby reducing the crack-driving force and increasing the fatigue resistance, was responsible for the slower fatigue crack growth rate with increasing K in comparison with Al–Cu–Mg–Ag alloy. Therefore, even though the fatigue fracture surface in Paris regime of Al–Cu–Mg–Ag–Er alloy exhibits a more brittle nature, i.e., cleavage-like transgranular fracture, as suggested by the presence of planar slip shown in Fig. 8(b), the slower FCG rate was still revealed.
Furthermore, the typical fatigue fracture surfaces in high growth rate regime for both alloys are shown in Fig. 9. Fig. 9(a)–(d) shows some representative features of the fatigue fracture surfaces in high growth rate regime of both alloys to highlight the differences between the fatigue surfaces of the alloys. Overall morphology of the fatigue fracture present in Fig. 9(a) and (b) was rough revealing regions of locally ductile failure, corresponding high magnification images of these regions are presented in Fig. 9(c) and (d). It can be noted that the region consisted predominantly of a dispersion of coarse dispersoid particles embedded within the microscopic voids in Al–Cu–Mg–Ag alloy, favored formation of crack at the interface of matrix and the coarse particles, resulting in a relative low ductility fracture. Tear ridges were also observed. Nevertheless, with Er addition, the pockets of fine shallow dimples distributed randomly throughout the surface, coupled with the coalescence of the fine microscopic voids resulting in fine microscopic cracks, suggests a great ductility fracture mechanism. Also, the absence or a great decrease of coarse dispersoid particles in the fracture surface (Fig. 9(d)) which favors a slight homogeneous deformation during cyclic loading, partly impedes the nucleation and coalescence of voids and then restrains the propagation of the microscopic cracks, eventually resulting in the low FCG rate. 4. Discussion Fatigue crack propagation behaviors of both underaged alloys shown great discrepancy, especially in Paris regime. Considering the fatigue fracture surfaces and the microstructures of both alloys, the crack propagation mode of Er-containing alloy evolves from
Fig. 6. FCG rates, da/dN, as a function of K for underaged Al–Cu–Mg–Ag–(Er) alloys.
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Fig. 7. Scanning electron micrographs of near-threshold fatigue fracture surfaces in underaged Al–Cu–Mg–Ag alloy (a), (c) and Al–Cu–Mg–Ag–Er alloy (b), (d) (crack growth direction from left to right).
ductile to crystallographic with the appearance of plenty of planar facets and secondary cracks. This means a close correlation between the fatigue crack propagation and the microstructure, although it is generally accepted that fatigue crack propagation in Paris regime is weakly dependent on the microstructure [25].
4.1. The influence of grain size The present of the persistent slip band (PSB) is beneficial to improve the fatigue propagation resistance by favoring the reversible cyclic slip and promoting the crack deflection [26]. The typical grain structures of both alloys used for FCP testing were given in Fig. 2, from which the grain sizes of Er-containing alloy were much larger than those of Al–Cu–Mg–Ag alloy. Accordingly, compared with Er-containing alloy, the finer grains of Al–Cu–Mg–Ag alloy were deleterious to the formation of PSB arising from the fact that the grain boundaries may act as barriers to PSB. For Al–Cu–Mg–Ag alloy as K increases to a high level in Paris regime, the plastic zone at the tip of the crack is greater than the grain size and promotes duplex slip, thereby minimizing the microstructural effects on FCG. Conversely, in the case of Er-containing alloy with large grains the sensitivity of FCG to the underlying microstructure keeps relatively stable as K increases. Generally, RICC that occurs in underaged Al-based alloys has also been considered as being particularly microstructure dependent as proposed by Refs. [27–29]. Large grains with planar slip provided better resistance of fatigue crack propagation by enhanced crack closure levels associated with increased surface roughness and crack deflection from the principal crack growth direction. It seems that the improved fatigue behavior of Er-containing microstructure arising from the high level of RICC. Actually though this crack closure mechanism may occur in the near-threshold regime to some extent, the influence of RICC is essentially negligible in Paris regime because the crack-tip opening displacements are of a size scale
greater than the average height of the fracture surface asperities as stress ratio R equals to 0.1.
4.2. Precipitates and secondary cracks The transition from the near-threshold regime to the Paris regime of fatigue crack growth is also accompanied by a noticeable change from a microstructure-sensitive to a microstructureinsensitive fracture behavior on condition that the average dimensions of the plastic zone at crack-tip is greater than the grain size [25]. This statement is fit for underaged Al–Cu–Mg–Ag alloy owing to its small grains. Nevertheless, considering the grain size scales of aged Er-containing alloy span from 0.5 mm to 1.2 mm and the plastic zone at crack-tip is relatively small, it implies the enhanced fatigue properties in Paris regime are also related to the precipitates ( and phases) within grains during aging treatment. For Al–Cu–Mg–Ag-based alloys, although it remains a topic of consideration discussion [30], the existence of shearing of the precipitates by gliding dislocation can be confirmed unambiguously as demonstrated in Ref. [31]. For both underaged microstructures, dislocations shear small semi-coherent precipitates to localize in planar bands of varying coarseness and glide reversibly during cyclic loading implying the high intrinsic FCG resistance. This planar-reversible slip improves the FCG resistance by favoring increasing crack closure and tip deflection [32]. In spite of the existence of planar-reversible slip, for Al–Cu–Mg–Ag alloys owing to the microstructure effects at high crack growth rates are limited, the gradual accumulation of cyclic micro-plastic deformation around the dispersoid particles, coupled with the ductile tearing in restricted areas, culminates in the nucleation of the microscopic cracks and ultimately leads to the fatigue damage. phase forms on {1 1 1}␣ habit planes and serves as an effective obstacle to dislocation gliding during deformation. Compared with the phase, the FCG resistance may be slightly degraded by the
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Fig. 8. Scanning electron micrographs of the fatigue fracture surfaces in Paris regime of underaged Al–Cu–Mg–Ag alloy (a), (c), (e) and Al–Cu–Mg–Ag–Er alloy (b), (d) and (f) (crack growth from left to right).
existence of the phase on {0 0 1}␣ planes. In the Er-containing microstructure, the presence of secondary cracks is apparent in Fig. 8. The Er-free and Er-containing microstructures, in conjunction with corresponding fatigue fracture surfaces of Paris regime, indicate that the crystallographic nature of fatigue secondary cracks is associated with the semi-coherent phase present in Fig. 3(d). That is, the formation of secondary cracks may due to the dislocations localized in the ␣/ -interphase boundaries that lead to the plastic deformation during cyclic loading. Since the phase forms as plates on {0 0 1}␣ planes, it is obvious that the propagation direction of secondary cracks is closed to {0 0 1}␣ planes and perpendicular to the principal crack as seen in Fig. 8. Furthermore, emphasis was placed on the evolution process of secondary cracks during cyclic loading. Details were shown in Figs. 8(f), from which the open and closed secondary cracks were clearly present. These secondary cracks with different morphology indicated that the crack closure happens inevitably during the FCP testing, although the level of this crack closure is limited. Secondary cracks propagation also occurs as arrowed in Fig. 8(f). Besides the planar-reversible slip, owing to the significant differences of fatigue fracture surfaces observed in both alloys, the crystallographic nature of secondary
cracks branched from the principal crack acts as a major extrinsic factor in decreasing the driving force for the growth of the fatigue crack that counteracts the deleterious effect of the phase sufficiently, hence improves the fatigue propagation resistance. 4.3. Mechanism governed by Er It is obvious that the addition of Er influences the precipitation sequence of Al–Cu–Mg–Ag alloy as seen in Fig. 3(a) and (b). Early stage decomposition at moderate temperature from the supersaturated solid solution of Al–Cu–Mg–Ag alloy leads to the formation of fine solute-rich zones, Mg-Ag co-clusters, which may serve as effective nucleation sites for the precipitation of the phase. Despite Mg and Ag atoms play an important role in the formation of the phase, here it is emphasized that the vacancies formed during quenching is believed to be a critical component for the precipitation of the phase because vacancies are indispensable to the diffusion of solute atoms. Diffusion of Mg and Ag atoms is realized by vacancy-solute exchanging mechanism and greatly depends on the movability of vacancy in Al matrix. As proposed by Hirosawa et al. [33], the characteristic behavior of Ag entirely arises from
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Fig. 9. Scanning electron micrographs of the fatigue fracture surfaces in high growth rate regime of underaged Al–Cu–Mg–Ag alloy (a), (c) and Al–Cu–Mg–Ag–Er alloy (b), (d) (crack growth from left to right).
both the strong interaction with Mg atoms and medium interaction with vacancies, suggesting that the formation and co-clusters stability are dependent on the vacancy concentration of the alloy and thereby Mg/Ag/Vacancy complexes should be responsible for the accelerated precipitation of the phase. According to the Al–Er phase diagram [34], the solubility of Er is small at room temperature in aluminum alloys. TEM and 3DAP observations confirmed that no Er-containing compounds were revealed within grains. Similarly, BSE image and Er mapping also exhibited the random distribution of Er. Er atoms, irrespective the existence of the Al8 Cu4 Er phase segregated at grain boundaries, should be partly soluted in the Al matrix. Owing to the large atom radius of Er, lattice distortion is induced which influence the distribution of vacancies. Despite the lack of the thermodynamic data about Er–vacancy binding energy, Er atoms may possess a strong binding with vacancies and exhibit the vacancy trap effect by collecting vacancies nearby in order to minimize the surrounding elastic strain. During initial aging at 165 ◦ C for 2 h, the formation of Mg–Ag co-clusters was disturbed on account of the lack of vacancies, thereby suppressing the precipitation of phase. As aging goes, the movability of vacancy was improved and the precipitation of phase was accelerated. Hence it is reasonable to speculate that vacancies, trapped during quenching from the solution treatment temperature, are partly if not most captured by Er atoms that are soluble in Al matrix during the initial artificial aging. This process, or an as-yet unresolved change, should be the reason for suppressing the formation of the Mg–Ag co-clusters as implied in Figs. 3 and 4. Also, during the initial aging, and precipitates are competing for copper. With increasing aging time the phase dissolution is suggested to be related to the loss of copper atoms arising from the precipitation and growth of phase in Al–Cu–Mg–Ag alloy as shown in Fig. 3(a) and (c), which is consistent with other observations [22]. However, the formation of precipitates retarded with Er content as illustrated in Fig. 3(b), except for the shortage of vacancies trapped by Er atoms mentioned
above, can partly be attributed to the precipitation kinetics of phase is slowed down by the formation of ternary Al8 Cu4 Er phase as it consumes copper atoms. Besides, it is proposed that the addition of Si also has a negative influence on the precipitation and stability of phases in Al–Cu–Mg–Ag alloy as the Mg/Si ratio (in terms of wt.%) is below 2.0 [35]. This is attributed to the strong Mg/Si interaction that disturbs the formation of pre-precipitate clusters which act as the nucleation sites for phase [36]. In the studies mentioned above, the composition of Si is relatively high. In this work, Si is treated as a trace element with a concentration on the order of 0.01 wt.% and Mg/Si ratio is far above 2.0. The deleterious effect of Si on the formation is negligible. Consequently, it is concluded that the excellent fatigue resistance of the Er-containing microstructure is attributed to the superior intrinsic nature of the microstructure.
5. Conclusions The influence of rare earth Er on the microstructures and fatigue fracture behavior of underaged Al–Cu–Mg–Ag alloy was investigated and could be summarized as follows:
(1) The dendritic substructure of as-cast Al–Cu–Mg–Ag alloy was remarkably refined by Er addition. During initial aging, the precipitation of phase was suppressed by the addition of Er, whereas the formation of phase was promoted. (2) The enhanced FCG resistance obtained in underaged Al–Cu–Mg–Ag–Er alloy indicates the fatigue crack growth is microstructural dependent. Large grain size of Er-containing alloy, accompanied with the presence of crystallographic secondary cracks, plays a dominant role in improving the FCG resistance of the studied alloy.
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Microstructures and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy with addition of rare earth Er Song Bai a,b , Zhiyi Liu a,b,∗ , Yuntao Li a,b , Yanhui Hou a,b,c , Xu Chen a,b a b c
Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Central South University, Changsha 410083, People’s Republic of China School of Material Science and Engineering, Central South University, Changsha 410083, People’s Republic of China School of Mechanics and Engineering, Southwest Jiaotong University, Chengdu 610031, People’s Republic of China
a r t i c l e
i n f o
Article history: Received 17 March 2009 Received in revised form 31 October 2009 Accepted 4 November 2009
Keywords: Al–Cu–Mg–Ag alloy Microstructures Fatigue fracture behavior Er
a b s t r a c t The effects of rare earth erbium (Er) on microstructures and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy were investigated. Microstructural examinations first revealed that the precipitation kinetics of phase was distinctly retarded by promoting the formation of phase with the Er addition during the initial aging. The fatigue crack propagation resistance of Er-containing microstructure was significantly enhanced arising from the presence of the crystallographic secondary cracks, which was directly relative to the large grain size. Results also suggested that the dendritic substructure of as-cast Al–Cu–Mg–Ag alloy was refined remarkably by Er addition. © 2009 Elsevier B.V. All rights reserved.
1. Introduction The addition of Ag with small amount is well known to enhanced the aging hardening response in a wide range of Al–Cu–Mg alloys by the precipitation of the fine and uniform plates termed phase on the {1 1 1}␣ planes [1,2]. phase is believed to be a variant of equilibrium phase (Al2 Cu) and keeps stable at elevated temperature arising from the high coarsening resistance [3]. In addition to the attractive strength properties [4], compared with commercial aluminium alloys, higher creep resistance was also exhibited if the alloys are in the underaged conditions [5–8]. These above make Al–Cu–Mg–Ag-based alloys the promising materials for aerospace applications. With the development of the aerospace industries, mechanical properties of Al-based alloys are required to be further improved to satisfy the applications that used in tension-dominated aircrafts require excellent fatigue damage tolerance. Furthermore, it is well accepted that small additions of rare earth elements are of great importance to improve the mechanical properties of Al-based alloys and provide extended alloying capabilities. Over the past several years, in the case of Al-based alloys rare earth Scandium (Sc) has been most commonly studied as beneficial microalloying element
∗ Corresponding author at: School of Material Science and Engineering, Central South University, Lu Mountain South Road, Hunan, Changsha 410083, People’s Republic of China. Tel.: +86 731 88836011; fax: +86 731 88876692. E-mail address: [email protected] (Z. Liu). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.11.011
[9–11], suggesting the beneficial effects on mechanical properties that associated with remarkably grain refinement, dynamic recrystallization controlling and also fatigue as proposed by Wirtz et al. [12]. However, the applications of Sc-containing alloys are extremely restricted due to the high cost of Sc addition. Instead, much attention was focused on rare earth Erbium (Er), which is much cheaper than Sc. A previous work conducted by Xu et al. [13] revealed that, with the addition of Er, the mechanical properties of Al–Mg alloys were strongly improved by the presence of the fine dispersed Al3 Er particles after heat treatment, whereas no similar effects were observed in a hot-extruded Al–Mg alloy without solution treatment as suggested by Wu et al. [14]. Based on the research by Karnesky et al. [15], the volume fraction of precipitates in the Al–0.08 at.% Sc alloy increased with small addition of 0.02 at.% Er. In Al–Zn–Mg alloys, Xu et al. [16] demonstrated that Er improves the strength considerably by precipitation strengthening mechanisms and grain refinement and retards recrystallization due to the pinning effect of fine dispersed A13 Er precipitates on dislocations and subgrain boundaries. Li et al. [17] investigated the influences of Er on the grain refinement in an as-cast Al–Cu–Mg–Ag alloy and found that Er atoms may segregate at grain boundary during solidification and resulted in ternary Al8 Cu4 Er phase. Besides, rare earth Cerium (Ce) and Ytterbium (Yb) also attract much attention. The addition of Ce promotes the precipitation and improves the thermal stability of the phase in Al–Cu–Mg–Ag alloy [18,19]. Similar results were also observed with addition of Yb [20]. Studies of the fatigue behavior of the Al–Cu–Mg–Ag-based alloys are limited [21], especially in the case of microalloying with rare
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Table 2 Room temperature mechanical properties of underaged Al–Cu–Mg–Ag–(Er) alloys.
Table 1 Compositions of the studied Al–Cu–Mg–Ag–(Er) alloys (in wt.%). Cu Er-free Er-containing
6.45 6.21
Mg 0.65 0.62
Ag 0.46 0.44
Mn 0.29 0.28
Ti 0.09 0.09
Zr 0.14 0.15
Er – 0.32
Fe 0.04 0.05
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Si 0.01 0.01
Al Bal. Bal.
earth Er. The present work mainly aims at clarifying the influence of rare earth element Er on the precipitation evolution and fatigue fracture behavior of an Al–Cu–Mg–Ag alloy during the initial aging. The mechanism governed by Er was also proposed from the obtained results.
Alloys
Temperature ◦
b (MPa)
0.2 (MPa)
ı (%)
Er-free
165 C × 2 h 165 ◦ C × 4 h
486.1 500.6
422.6 476.4
16.2 11.3
Er-containing
165 ◦ C × 2 h 165 ◦ C × 6 h
453.7 488.2
341.9 460.1
21.0 11.7
condition), the FCP data was considered as a reference for further design. 2.3. Microstructural analysis
2. Experimental procedures 2.1. Sample preparation and heat treatment Al–Cu–Mg–Ag alloys with high Cu/Mg ratios are used in this study and have the compositions as listed in Table 1. It should be noted that although the relative amount of Cu in Al–Cu–Mg–Ag–Er alloy is a little lower as compared to Al–Cu–Mg–Ag alloy, it has no significant influence on the experimental results. Both alloys were cast into a cast iron mould. The ingots of Al–Cu–Mg–Ag–(Er) alloys experienced homogenization followed by air cooling to room temperature, scalped and hot roll to 2.2 mm thick strips at about 450 ◦ C. The alloy samples were solution treated at 515 ◦ C for 6 h, cold water quenched and immediately aged at 165 ◦ C in air furnace for various times. 2.2. Tensile and FCP testing Specimens for tensile testing were prepared vertical to the longitudinal direction of the strips with gauge length 30 mm and 2.0 mm thick. Tensile testing was performed at room temperature on the CSS-44100 type testing machine with 2 mm/min loading speed. Fatigue crack propagation (FCP) testing was performed on compact tension (CT) specimens taken from the strips in the L–T orientation with a size (in mm) of 47.5 × 45 × 2 (L × W × B) to obtain the fatigue crack growth (FCG) rates that are above 10−5 mm/cycle. The underaged samples were chosen for FCP testing. All FCP tests were conducted at a stress ratio (R = Kmin /Kmax ) of 0.1 with a loading frequency of 10 Hz on an Instron 8801 50 kV fatigue tester at room temperature, laboratory air environment. The crack length was monitored by elastic compliance technique. Though the plane stress condition was inevitably present during FCP testing and not in accordance with the standard of fatigue test, here it is necessary to pointed out that the purpose of this study was to clarify the fatigue behavior of Er-containing alloy in thin sheet which is closest to its condition in real applications (that is plane stress
Optical metallographic examination was used to reveal the grain structures of both alloys. Transmission electron microscopy (TEM), along with selected area electron diffraction (SAED), was employed to characterize the differences of the microstructure evolution in Al–Cu–Mg–Ag alloys with and without Er during aging treatment. Samples suitable for TEM were thin slices with 3 mm in diameter and electro-polished by using twin-jet equipment with a voltage of 10 V in a 70% methanol and 30% nitric acid at approximately −20 ◦ C, after which the slices were cleaned in ethanol at room temperature for at least 5 min. These were then examined on a Tecnai G2 20 ST TEM machine operating at 200 kV. Fatigue fracture surfaces of the samples cyclically deformed in FCP testing was analyzed by KYKY 2800 scanning electron microscopy (SEM) and the distribution of rare earth Er was measured by Quanta 200 machine. Three-dimensional atom probe (3DAP) technique was also used to reveal how the addition of Er influences the precipitation reactions in Al–Cu–Mg–Ag alloy. 3. Results 3.1. Microstructures and mechanical properties The typical grain structures of the as-cast alloys used in this study were characterized by optical microscopy in order to reveal the influence of Er addition as shown in Fig. 1. Fig. 1(a) illustrates the cast structure of Al–Cu–Mg–Ag alloy that predominated by large equiaxed grains with coarsening dendritic substructure. The grain boundaries were decorated by continuous phases. Great branch distance was present and the average grain size of this casting was about 0.4 mm. Apparently, the distance between the branches of the dendritic substructure was greatly reduced as a direct consequence of the addition of Er shown in Fig. 1(b), which is consistent with earlier observations [17] on similar materials. The grain was slightly refined and was about 0.2 mm. Grain structures of the underaged samples used in the FCP testing were shown in Fig. 2,
Fig. 1. Typical grain structures of as-cast alloys (a) Al–Cu–Mg–Ag and (b) Al–Cu–Mg–Ag–Er.
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Fig. 2. Typical grain structures of the FCP testing samples (a) Al–Cu–Mg–Ag alloy aged at 165 ◦ C for 4 h and (b) Al–Cu–Mg–Ag–Er alloy aged at 165 ◦ C for 6 h.
which both show the fully recrystallised grains. The Er-containing alloy exhibits a much larger recrystallised grain structure. The microstructures of underaged Al–Cu–Mg–Ag–(Er) alloys were characterized as shown in Fig. 3, corresponding room temperature mechanical properties are given in Table 2. The microstructure of Al–Cu–Mg–Ag alloy dominated by phase formed on the {1 1 1}␣ planes, together with fine plates on the {0 0 1}␣ planes, was found after aging for 2 h at 165 ◦ C as shown in Fig. 3(a), which was consistent with SAED pattern. Aging for 4 h at 165 ◦ C mostly produced phase and phase was only present as a minor phase as seen in Fig. 3(c), indicating the precipitation of phase was enhanced as aging goes at the expense of phase supported by previous study [22]. Corresponding SAED pattern in Fig. 3(c) was hard to reveal the spots that associated with phase, only reflections at the 1/3 and 2/3 {0 2 2}␣ positions produced by phase were observed. However, as shown in Fig. 3(b), with the addition of Er differences were distinctly observed in the appearance of phase compared with Fig. 3(a). Only phase was present and close examination of the microstructure in the thin TEM samples taken near 1 1 0␣ zone axis failed to reveal any precipitates that could be identified as phase. SAED pattern taken near 0 0 1␣ zone axis in Fig. 3(b) only confirmed the existence of phase. Considering the competition between and phases, it is suggested that the role of rare earth Er is to suppress the precipitation of phase by promoting the formation of phase at early stage of aging. The precipitation of phase was considerably accelerated following aged at 165 ◦ C for 6 h as revealed by the fine and uniform dispersions of phase along with plenty of fine precipitates observed in Fig. 3(d) and corresponding SAED pattern shown obvious reflections of and precipitates, although the intensity of phase was extremely weak, notably that the redissolving of phase occurred in Al–Cu–Mg–Ag alloy (seen in Fig. 3(c)) have been remarkably impeded. It is generally accepted that, in Al–Cu–Mg–Ag alloy, the precipitate kinetics of phase precipitates is greatly accelerated by Ag addition, since Ag and Mg atoms are strongly segregated to the broad face of phase as aging goes, lying on the {1 1 1}␣ planes, which reduces the misfit strain at the ␣/ interface and promotes the growth of phase [2,22,23]. Consequently, 3DAP was employed to clarify the microalloying behavior of Er addition. Despite the limitation of the field range, however, 3DAP result failed to reveal any evidence of Er atoms segregate to the ␣/ interface as seen in Fig. 4. This may attribute to the trace element of Er in this study.
Although Er atoms in as-cast Al–Cu–Mg–Ag alloy mainly segregate at grain boundary during solidification and resulted in ternary Al8 Cu4 Er phase [17], part of which are soluble in Al matrix during heat treatment, if not, the retardant of precipitation of phase is hard to explain because the Al8 Cu4 Er phase has no direct effect on the precipitation in Al matrix. Fig. 5 illustrates the backscattered electron (BSE) image of underaged Al–Cu–Mg–Ag–Er alloy (165 ◦ C × 6 h) and elements mapping of Cu and Er. It was clear that coarse Al8 Cu4 Er phase in as-cast structure was not visible and the distribution of Cu and Er was random. No obvious co-segregation of Cu and Er was present. This implies that Er atoms soluted in the Al matrix are associated with the slow precipitation response at the initial aging. 3.2. Fatigue fracture behavior As seen in Table 2 two studied underaged conditions (165 ◦ C × 4 h of Al–Cu–Mg–Ag alloy and 165 ◦ C × 6 h of Al–Cu–Mg–Ag–Er alloy) show very similar values of room temperature mechanical properties. Therefore, it is expected that different fatigue crack growth rates of two alloys are not resulted, at least, from the differences of mechanical properties. The variation of FCG rates with the stress intensity factor range (K) is given in Fig. 6 for the comparison between both underaged alloys. Apparently, both da/dN–K curves show the three typical stages from Kth up to the critical stress intensity amplitude where the final fatigue fracture occurs. Contrary to the expected above, even though no pronounced differences in FCG rates exist in nearthreshold regime (below K values of about 13 MPa m0.5 ), some clear deviation between the two curves was observed in the intermediate K region, Paris regime, where the Al–Cu–Mg–Ag alloy showed high FCG rate whereas higher fatigue crack propagation resistance against slow FCG rate was revealed for the microstructure of Al–Cu–Mg–Ag–Er alloy. It also should be emphasized that even at high K region (above K values of about 28 MPa m0.5 ), comparing with Al–Cu–Mg–Ag alloy which was obviously seen that FCG rate accelerated dramatically when K equals to 28 MPa m0.5 and then followed by fast fatigue fracture, the Er-containing alloy still exhibited higher fatigue crack propagation resistance and was insensitive to the variety of K, i.e., the FCG rate of Al–Cu–Mg–Ag alloy was several time higher than that of Er-containing alloy at K of 30 MPa m0.5 . The fatigue fracture of underaged Al–Cu–Mg–Ag–Er alloy occurred ultimately at about K of 42 MPa m0.5 . As revealed in Fig. 6, the crack propagation resistance was greatly enhanced with
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Fig. 3. TEM bright field micrographs and corresponding SAED patterns of the underaged Al–Cu–Mg–Ag alloy aged for (a) 2 h and (c) 4 h and Al–Cu–Mg–Ag–Er alloy aged for (b) 2 h and (d) 6 h at 165 ◦ C. Electron beam is approximately parallel to 1 1 0␣ zone axis in (a), (c) and (d) versus 0 0 1␣ zone axis in (b).
the addition of Er, especially in the high K values. It is assumed that the role of Er addition on the fatigue fracture behavior of the Al–Cu–Mg–Ag alloy should be concluded to stabilize the fatigue crack propagation as K increases. Examination of the fatigue fracture surfaces of both underaged alloys were carried out for comparisons at (i) low magnification to clarify the overall fatigue fracture morphology of different regimes and (ii) high magnification to identify the characteristics of the fatigue cracks growth and the fine-scale fatigue fracture features for both alloys in FCP testing. Representative fatigue fracture features were shown in Figs. 7–9 taken from fatigue fracture regions where the fatigue crack growth rates were measured. The near-threshold fatigue fracture surfaces of underaged alloys given in Fig. 7 shows different morphology. Compared with Fig. 7(a) and (b), the surface of Al–Cu–Mg–Ag alloy was characterized by minor fibrous beach marks that are roughly parallel to the prin-
cipal fatigue crack growth direction versus many main shearing ridges with the addition of Er were observed. According to the high magnification view shown in Fig. 7(c) and (d), the presence of microscopic voids arising from the localized micro-plastic deformation at the incoherent Mn/Zr-rich dispersoid particles through the matrix in Al–Cu–Mg–Ag alloy were found, which were not present in Er-containing alloy. It is generally accepted that the voids can be considered as being equivalent to second-phase particles having zero stiffness and facilitate a gradual increase of both the size and relative volume fraction of voids in cyclic deformation that ultimately result in the increments of crack-tip extension [24]. However, with the addition of Er, the fracture surface seems to be comprised of crystallographic shearing ridges joined by each other as revealed in Fig. 7(d). Crack deviation occurs apparently because the crack follows these weakened planes and deviates 20–40◦ with respect to the principal crack growth direction. FCG in this region
Fig. 4. 3DAP elemental mapping of a selected region of the Al–Cu–Mg–Ag–Er alloy aged for 6 h at 165 ◦ C.
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Fig. 5. BSE image and elements mapping of Al–Cu–Mg–Ag–Er alloy aged at 165 ◦ C for 6 h.
occurs by a transgranular shear mechanism. In this study, at low K values although the fatigue fracture surfaces did shown some differences as seen in Fig. 7, the FCG rates were almost the same for both alloys shown in Fig. 6. As shown in Fig. 8, it was observed that significant differences occurred in Paris regime between underaged Al–Cu–Mg–Ag and Al–Cu–Mg–Ag–Er alloys for which the most distinct deviation in FCG rates were measured as revealed in Fig. 6. The overall morphology of the fatigue fracture surfaces in Paris regime for both alloys was shown in Fig. 8(a) and (b), both exhibiting fairly well-developed fatigue characteristics and corresponding high magnification observations were present in Fig. 8(c)–(f), respectively. In Al–Cu–Mg–Ag alloy, the region of stable fatigue crack propagation comprised of pockets of fibrous shell marks formed during cyclic deformation and widely dispersed microscopic voids formed around dispersoid particles were observed. In contrast, few distinct microscopic voids were shown with the addition of Er. Fracture was predominantly transgranular with a random distribution of slip steps on which many regular coplanar cracks, termed secondary cracks, formed. Evidence of some ductile tearing arrowed in Fig. 8(b) was also exhibited. It can be clearly seen, by comparison between Fig. 8(c) and (d), that the transgranular fracture regions of Al–Cu–Mg–Ag alloy was predominantly comprised of a dispersion of microscopic voids and spherical shaped Mn/Zr-rich dispersoid particles of varying sizes, although some microscopic cracks were present as arrowed and circled in Fig. 8(c), whereas the presence of secondary cracks in a much higher density was visible in Er-containing alloy. Furthermore, the orientation of secondary cracks alignment was not along the principal fatigue crack growth direction but normal to the fracture surface. Fig. 8(e) and (f) illustrated some details of the fatigue fracture regions circled in Fig. 8(c) and (d), respectively. Secondary cracks were practically absent in Al–Cu–Mg–Ag alloy, only microscopic cracks were observed as arrowed in Fig. 8(e). Fig. 8(f) exhibits the presence of arrays of secondary cracks in Er-containing alloy and it was also obvious that minor uniformly spaced fatigue striations appear as arrowed in the surface of secondary cracks, which directly indicate the propagation process of secondary cracks also occurs during the cyclic deformation. SEM observations analysis revealed that the occurrence of secondary cracks results in a more complex crack path, thereby reducing the crack-driving force and increasing the fatigue resistance, was responsible for the slower fatigue crack growth rate with increasing K in comparison with Al–Cu–Mg–Ag alloy. Therefore, even though the fatigue fracture surface in Paris regime of Al–Cu–Mg–Ag–Er alloy exhibits a more brittle nature, i.e., cleavage-like transgranular fracture, as suggested by the presence of planar slip shown in Fig. 8(b), the slower FCG rate was still revealed.
Furthermore, the typical fatigue fracture surfaces in high growth rate regime for both alloys are shown in Fig. 9. Fig. 9(a)–(d) shows some representative features of the fatigue fracture surfaces in high growth rate regime of both alloys to highlight the differences between the fatigue surfaces of the alloys. Overall morphology of the fatigue fracture present in Fig. 9(a) and (b) was rough revealing regions of locally ductile failure, corresponding high magnification images of these regions are presented in Fig. 9(c) and (d). It can be noted that the region consisted predominantly of a dispersion of coarse dispersoid particles embedded within the microscopic voids in Al–Cu–Mg–Ag alloy, favored formation of crack at the interface of matrix and the coarse particles, resulting in a relative low ductility fracture. Tear ridges were also observed. Nevertheless, with Er addition, the pockets of fine shallow dimples distributed randomly throughout the surface, coupled with the coalescence of the fine microscopic voids resulting in fine microscopic cracks, suggests a great ductility fracture mechanism. Also, the absence or a great decrease of coarse dispersoid particles in the fracture surface (Fig. 9(d)) which favors a slight homogeneous deformation during cyclic loading, partly impedes the nucleation and coalescence of voids and then restrains the propagation of the microscopic cracks, eventually resulting in the low FCG rate. 4. Discussion Fatigue crack propagation behaviors of both underaged alloys shown great discrepancy, especially in Paris regime. Considering the fatigue fracture surfaces and the microstructures of both alloys, the crack propagation mode of Er-containing alloy evolves from
Fig. 6. FCG rates, da/dN, as a function of K for underaged Al–Cu–Mg–Ag–(Er) alloys.
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Fig. 7. Scanning electron micrographs of near-threshold fatigue fracture surfaces in underaged Al–Cu–Mg–Ag alloy (a), (c) and Al–Cu–Mg–Ag–Er alloy (b), (d) (crack growth direction from left to right).
ductile to crystallographic with the appearance of plenty of planar facets and secondary cracks. This means a close correlation between the fatigue crack propagation and the microstructure, although it is generally accepted that fatigue crack propagation in Paris regime is weakly dependent on the microstructure [25].
4.1. The influence of grain size The present of the persistent slip band (PSB) is beneficial to improve the fatigue propagation resistance by favoring the reversible cyclic slip and promoting the crack deflection [26]. The typical grain structures of both alloys used for FCP testing were given in Fig. 2, from which the grain sizes of Er-containing alloy were much larger than those of Al–Cu–Mg–Ag alloy. Accordingly, compared with Er-containing alloy, the finer grains of Al–Cu–Mg–Ag alloy were deleterious to the formation of PSB arising from the fact that the grain boundaries may act as barriers to PSB. For Al–Cu–Mg–Ag alloy as K increases to a high level in Paris regime, the plastic zone at the tip of the crack is greater than the grain size and promotes duplex slip, thereby minimizing the microstructural effects on FCG. Conversely, in the case of Er-containing alloy with large grains the sensitivity of FCG to the underlying microstructure keeps relatively stable as K increases. Generally, RICC that occurs in underaged Al-based alloys has also been considered as being particularly microstructure dependent as proposed by Refs. [27–29]. Large grains with planar slip provided better resistance of fatigue crack propagation by enhanced crack closure levels associated with increased surface roughness and crack deflection from the principal crack growth direction. It seems that the improved fatigue behavior of Er-containing microstructure arising from the high level of RICC. Actually though this crack closure mechanism may occur in the near-threshold regime to some extent, the influence of RICC is essentially negligible in Paris regime because the crack-tip opening displacements are of a size scale
greater than the average height of the fracture surface asperities as stress ratio R equals to 0.1.
4.2. Precipitates and secondary cracks The transition from the near-threshold regime to the Paris regime of fatigue crack growth is also accompanied by a noticeable change from a microstructure-sensitive to a microstructureinsensitive fracture behavior on condition that the average dimensions of the plastic zone at crack-tip is greater than the grain size [25]. This statement is fit for underaged Al–Cu–Mg–Ag alloy owing to its small grains. Nevertheless, considering the grain size scales of aged Er-containing alloy span from 0.5 mm to 1.2 mm and the plastic zone at crack-tip is relatively small, it implies the enhanced fatigue properties in Paris regime are also related to the precipitates ( and phases) within grains during aging treatment. For Al–Cu–Mg–Ag-based alloys, although it remains a topic of consideration discussion [30], the existence of shearing of the precipitates by gliding dislocation can be confirmed unambiguously as demonstrated in Ref. [31]. For both underaged microstructures, dislocations shear small semi-coherent precipitates to localize in planar bands of varying coarseness and glide reversibly during cyclic loading implying the high intrinsic FCG resistance. This planar-reversible slip improves the FCG resistance by favoring increasing crack closure and tip deflection [32]. In spite of the existence of planar-reversible slip, for Al–Cu–Mg–Ag alloys owing to the microstructure effects at high crack growth rates are limited, the gradual accumulation of cyclic micro-plastic deformation around the dispersoid particles, coupled with the ductile tearing in restricted areas, culminates in the nucleation of the microscopic cracks and ultimately leads to the fatigue damage. phase forms on {1 1 1}␣ habit planes and serves as an effective obstacle to dislocation gliding during deformation. Compared with the phase, the FCG resistance may be slightly degraded by the
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Fig. 8. Scanning electron micrographs of the fatigue fracture surfaces in Paris regime of underaged Al–Cu–Mg–Ag alloy (a), (c), (e) and Al–Cu–Mg–Ag–Er alloy (b), (d) and (f) (crack growth from left to right).
existence of the phase on {0 0 1}␣ planes. In the Er-containing microstructure, the presence of secondary cracks is apparent in Fig. 8. The Er-free and Er-containing microstructures, in conjunction with corresponding fatigue fracture surfaces of Paris regime, indicate that the crystallographic nature of fatigue secondary cracks is associated with the semi-coherent phase present in Fig. 3(d). That is, the formation of secondary cracks may due to the dislocations localized in the ␣/ -interphase boundaries that lead to the plastic deformation during cyclic loading. Since the phase forms as plates on {0 0 1}␣ planes, it is obvious that the propagation direction of secondary cracks is closed to {0 0 1}␣ planes and perpendicular to the principal crack as seen in Fig. 8. Furthermore, emphasis was placed on the evolution process of secondary cracks during cyclic loading. Details were shown in Figs. 8(f), from which the open and closed secondary cracks were clearly present. These secondary cracks with different morphology indicated that the crack closure happens inevitably during the FCP testing, although the level of this crack closure is limited. Secondary cracks propagation also occurs as arrowed in Fig. 8(f). Besides the planar-reversible slip, owing to the significant differences of fatigue fracture surfaces observed in both alloys, the crystallographic nature of secondary
cracks branched from the principal crack acts as a major extrinsic factor in decreasing the driving force for the growth of the fatigue crack that counteracts the deleterious effect of the phase sufficiently, hence improves the fatigue propagation resistance. 4.3. Mechanism governed by Er It is obvious that the addition of Er influences the precipitation sequence of Al–Cu–Mg–Ag alloy as seen in Fig. 3(a) and (b). Early stage decomposition at moderate temperature from the supersaturated solid solution of Al–Cu–Mg–Ag alloy leads to the formation of fine solute-rich zones, Mg-Ag co-clusters, which may serve as effective nucleation sites for the precipitation of the phase. Despite Mg and Ag atoms play an important role in the formation of the phase, here it is emphasized that the vacancies formed during quenching is believed to be a critical component for the precipitation of the phase because vacancies are indispensable to the diffusion of solute atoms. Diffusion of Mg and Ag atoms is realized by vacancy-solute exchanging mechanism and greatly depends on the movability of vacancy in Al matrix. As proposed by Hirosawa et al. [33], the characteristic behavior of Ag entirely arises from
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Fig. 9. Scanning electron micrographs of the fatigue fracture surfaces in high growth rate regime of underaged Al–Cu–Mg–Ag alloy (a), (c) and Al–Cu–Mg–Ag–Er alloy (b), (d) (crack growth from left to right).
both the strong interaction with Mg atoms and medium interaction with vacancies, suggesting that the formation and co-clusters stability are dependent on the vacancy concentration of the alloy and thereby Mg/Ag/Vacancy complexes should be responsible for the accelerated precipitation of the phase. According to the Al–Er phase diagram [34], the solubility of Er is small at room temperature in aluminum alloys. TEM and 3DAP observations confirmed that no Er-containing compounds were revealed within grains. Similarly, BSE image and Er mapping also exhibited the random distribution of Er. Er atoms, irrespective the existence of the Al8 Cu4 Er phase segregated at grain boundaries, should be partly soluted in the Al matrix. Owing to the large atom radius of Er, lattice distortion is induced which influence the distribution of vacancies. Despite the lack of the thermodynamic data about Er–vacancy binding energy, Er atoms may possess a strong binding with vacancies and exhibit the vacancy trap effect by collecting vacancies nearby in order to minimize the surrounding elastic strain. During initial aging at 165 ◦ C for 2 h, the formation of Mg–Ag co-clusters was disturbed on account of the lack of vacancies, thereby suppressing the precipitation of phase. As aging goes, the movability of vacancy was improved and the precipitation of phase was accelerated. Hence it is reasonable to speculate that vacancies, trapped during quenching from the solution treatment temperature, are partly if not most captured by Er atoms that are soluble in Al matrix during the initial artificial aging. This process, or an as-yet unresolved change, should be the reason for suppressing the formation of the Mg–Ag co-clusters as implied in Figs. 3 and 4. Also, during the initial aging, and precipitates are competing for copper. With increasing aging time the phase dissolution is suggested to be related to the loss of copper atoms arising from the precipitation and growth of phase in Al–Cu–Mg–Ag alloy as shown in Fig. 3(a) and (c), which is consistent with other observations [22]. However, the formation of precipitates retarded with Er content as illustrated in Fig. 3(b), except for the shortage of vacancies trapped by Er atoms mentioned
above, can partly be attributed to the precipitation kinetics of phase is slowed down by the formation of ternary Al8 Cu4 Er phase as it consumes copper atoms. Besides, it is proposed that the addition of Si also has a negative influence on the precipitation and stability of phases in Al–Cu–Mg–Ag alloy as the Mg/Si ratio (in terms of wt.%) is below 2.0 [35]. This is attributed to the strong Mg/Si interaction that disturbs the formation of pre-precipitate clusters which act as the nucleation sites for phase [36]. In the studies mentioned above, the composition of Si is relatively high. In this work, Si is treated as a trace element with a concentration on the order of 0.01 wt.% and Mg/Si ratio is far above 2.0. The deleterious effect of Si on the formation is negligible. Consequently, it is concluded that the excellent fatigue resistance of the Er-containing microstructure is attributed to the superior intrinsic nature of the microstructure.
5. Conclusions The influence of rare earth Er on the microstructures and fatigue fracture behavior of underaged Al–Cu–Mg–Ag alloy was investigated and could be summarized as follows:
(1) The dendritic substructure of as-cast Al–Cu–Mg–Ag alloy was remarkably refined by Er addition. During initial aging, the precipitation of phase was suppressed by the addition of Er, whereas the formation of phase was promoted. (2) The enhanced FCG resistance obtained in underaged Al–Cu–Mg–Ag–Er alloy indicates the fatigue crack growth is microstructural dependent. Large grain size of Er-containing alloy, accompanied with the presence of crystallographic secondary cracks, plays a dominant role in improving the FCG resistance of the studied alloy.
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