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Strength and fracture of aluminum alloys
Materials Science and Engineering A280 (2000) 8 – 16 www.elsevier.com/locate/msea
Strength and fracture of aluminum alloys T. Kobayashi * Department of Production Systems Engineering, Toyohashi Uni6ersity of Technology, Tempaku, Toyohashi 4418580, Japan Received 14 June 1999
Abstract In aluminum alloys, it is general that dimple type fracture occurs from inclusions or second phases particles. Intergranular, local shear and delamination type fractures are also sometimes observed. On Al – Li alloy (A2091), effect of inclusion particles on fracture behavior is analyzed using HRR singularity and Eshelby type internal stress analysis. Fracture of inclusions of CuAl2 and Al2CuMg is observed (5–8 mm) and their strength is estimated about 710 MPa and decreased with increasing particle size. Fracture of Al3Zr or Al3Ti particles is not observed. Moreover, it is clarified that delamination occurs when the normal stress against the grain boundary attains 116 and 113 MPa at liquid He and room temperature, respectively. Moreover, newly developed high strength Al–Nb wire with 1 Gpa is introduced. Fracture and fatigue behaviors of aluminum casting alloys are also stated. Increasing the iron content, fracture toughness and fatigue properties (also impact fatigue properties) are degraded. The effect of Ca addition that possesses the modifying effect of the eutectic Si morphology is examined. It has been observed that the harmful effect of iron is largely improved. Increasing the Si content in Al – Si casting alloys, it is generally observed that fatigue strength or DKth is increased. Toughness of MMC is generally rather low. Based on the simulation analysis on fracture of Al–SiCW composite, high toughness MMC has been already developed, where some aggregated SiCW granules have been embedded into A6061 matrix. According to FEM analysis, it has been suggested that there is an appropriate aggregation ratio in the granule (i.e. SiCW content in the granule). The above topics have been reviewed. Published by Elsevier Science S.A. Keywords: Fracture toughness; Al–Li alloy; Aluminum casting alloy; Al – Nb wire, MMCs
1. Introduction The researches of strength and fracture of aluminum alloys have been actively carried out mainly in the aircraft industry related field. Generally ductile fractures occur with the sequence of nucleation, growth and coalescence of voids. It is very important to accurately understand the relationship between fracture behavior and microstructure for the purpose of alloy design. In aluminum alloys, generally, there is no ductile– brittle transition phenomenon observed in BCC metals. This provides a problem for improving the ductile fracture resistance when the yield strength is increased. Generally the toughness of the overaged alloys tends to decrease in comparison with the underaged ones. The unstable fast fracture becomes frequent, even if it is a ductile fracture, because the stengthening lowers the
* Tel.: +81-532-44-6693; fax: + 81-532-44-6690. E-mail address: [email protected] (T. Kobayashi) 0921-5093/00/$ - see front matter Published by Elsevier Science S.A. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 6 4 9 - 8
level of toughness. This becomes a problem with largescale structures. In addition, fracture characteristics under the impact load seem to become also important, because the application to automobiles and vehicles will increase. It is also predicted that the use of casting materials and aluminum matrix composites will also be expanded. It is, therefore, important to accumulate the data on the safety for the fracture behavior of aluminum alloys. In this review paper, it is reported on the problems of strength and fracture from the above viewpoints for wrought and casting aluminum alloys and composites, mainly based on the author’s studies.
2. Features of strength and fracture in aluminum alloys Since many second phase particles are contained, the fracture of aluminum alloys is the ductile fracture of the dimple formation type. In this case, whether it forms the dimple by fracture of the particle itself or by debonding at the particle interface, will change accord-
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
ing to its properties, shapes, volume fraction and condition of matrix interface of the particle. The outline of ductile fracture of void formation type is generally described in the following text. The ductile fracture progresses generally by the process of nucleation-growth-coalescence of the voids. Fig. 1 shows the various types of void in such case [1]. The nucleus of the void is inclusion, dispersoid or precipitate particle. In titanium alloys, such inclusion particles often do not exist. It is considered that in this case, void is nucleated by the mechanism of interference of grain boundary and slip band shown in Fig. 1(d) and (e). The void grows, when it is nucleated, the parameter on stress triaxiality sm /s¯ :sm is hydrostatic component of tensile stress, s is effective stress of sm) is very important. It is possible to obtain the stress triaxiality from the equation of Bridgeman in the tensile test [2], [3]. sm/s=1/3 + ln(a/2R +1)
(1)
where a is the smallest radius in void growth in the tensile test specimen, R is a projection radius of curvature. It is related to the stress triaxiality, Rice and Tracey give the following equation as Rv in respect of the void growth rate [4].
Fig. 1. Schematic illustration of void nucleation. Schwalbe et al. [1].
dRv/Rv = 0.28 dop exp(1.5sm/s)
(2)
where op is an effective plastic strain. Fig. 2 shows the results of examining this relation in 7075 system aluminum alloys [3]. It is clear that the ratio of void growth increases by increasing the stress triaxiality, and there is the rectilinear relation. It is also proven that coefficient of 0.28 differs by the alloys. It is shown that the void is easy to grow regardless of the existence of nucleus of the void, when the stress triaxiality increases. The coalescence of voids has not always been clarified, since it is unstably and rapidly occurred. McClintock [5] considered that adjoining voids grow and coalesce. On the other hand, the small voids may support coalescence shown in Fig. 1(b). From a past, however, the mechanism coupled by the shear by the local strain concentration (formation of void sheet) between voids has been indicated. Fig. 1(f) indicates that the straight line, which connects the void root, becomes p/4 rad (45°), when the void length (depth) becomes equal to a distance between voids [6]. Fig. 3 shows deformation models in the age-hardenable alloys [7]. In age-hardenable aluminum alloys, precipitate particles (0.02–0.1 mm) are sheared by the dislocations. By inducing the coarse slip over the grain size, this may cause the embrittlement (type A). In type B, by precipitating the particle and forming PFZ (precipitate free zone) in grain boundary vicinity, and coarse particles precipitate on grain boundary, the intergranular fracture is easy to be caused. In Al–Li alloy, the fracture patterns of A and B types are mainly observed. Especially, there is the high possibility of connecting with the macroscopic shear fracture in case of A type. However, when to some extent large dispersoid particles exist like type C, toughness is improved. In case of B type, the local stress concentration at PFZ increases when the difference of (sM –sPFZ) increases, where the strengths of matrix and PFZ define sM and sPFZ, respectively, the fracture tendency at PFZ increases. However, most large factor is void nucleation at coarse particle and after extension growth on grain boundary. For the fracture of void formation type like the above, Garrett et al. [8] give the following equation as a result of occurring the fracture, when the strain at a main crack-tip becomes threshold value, o*. c 2 2 1/2 KIC : (2CEo*·s c yn /1−6 )
Fig. 2. Relationship between void growth rate and stress triaxiality in 7075 aluminum alloys. (Zr content (mass%) AB 0.01, B:0.06, C:0.08, D:0.16, E:030)
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(3)
where C is constant (= 1/40), sy is yield stress, n is work hardening exponent and 6 is Poisson ratio. On the other hand, Ritchie et al. [9] have indicated that the strain at a main crack-tip must have exceeded limitative fracture strain o*(s c m/s) over a characteristic distance l*0 (order of process zone= CTOD) in case of such strain dominant fracture. In 3-point bending test on above-mentioned 7075 system aluminum alloy (E),
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T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
3. Wrought alloys
3.1. Al–Li alloys
Fig. 3. Deformation and fracture model of age-hardenable alloys. (Welpmann et al. [7]).
In Al–Li alloy, the age hardening occurs by the precipitate process below. SSSS G.P. zone d%(Al3Li)d(AlLi)
Fig. 4. Distribution of plastic strain and critical fracture strain as a function of stress triaxiality at a main crack-tip. X is the distance from a main crack tip, x1 is the first void nucleation point (7075 alloy).
Fig. 5. Variation of maximum load (Pm) and absorbed energy (Et) as a function of test temperature obtained by CAI system in 8090 aluminum alloy.
the example in search of ratio between distance from a main crack-tip, x and CTODIC (dIC) and relationship between local effective plastic strain, op is shown in Fig. 4. It has been confirmed that intersection point, xI between op and o*(s f m/s) corresponds to distance, l between second phase particles in the first void nucleation point [3]. The void width and depth are defined as W and h, and the roughness of fracture surface is defined as M =h/M, of = ln(M 2/3f)/3 and, JIC =s0ofl*0 can be considered ( f: volume fraction of particle,s0: flow stress), JIC :s0/3 ln(M 2/3f)·l*. 0
(4)
On the other hand, Hahn et al. [10] give following equation, the crack occurs, when the second phase particle comes in the process zone. KIC :[2syE(p/6)1/3d]1/2f − 1/6
(5)
where E is Young’s modulus, d is diameter of second phase particle. Though this equation agrees with the experimental value, the contradiction has also been indicated for KIC rising in proportion to d 1/2 and s 1/2 y .
The globular and coherent d% phase mainly contributes to the hardening. However, d% phase is easy to be cut off by dislocations, and the sedimentation of the dislocation to the grain boundary is easy to pile up. The local shearing deformation is easily frequent, and there is the high possibility of separation and intergranular fracture in the slip plane. In addition, in this alloy, the absorption of the hydrogen gas is frequent, and low toughness is a problem. However, this alloy shows the positive temperature dependence of the strength observed in Ni based superalloys, and it is excellent for the high temperature strength. Whereas, it is excellent in toughness under low temperature, and it has also been applied to the fuel tank of space shuttle. Fig. 5 shows the example on such enhancement of low-temperature toughness [11]. This is related to the delamination which originates from the fibrous structure by the rolling, and the effect of the anisotropy seems to be big [12]. Ritchie et al. [12] have asserted that such exogenous factor is a primary cause on the enhancement of low-temperature toughness [13,14]. Recently the author and co-workers have examined the initiation stress of delamination crack of the A2091 alloy (L-T direction) by the ultrasonic wave measurement at room temperature and liquid helium temperature [15]. According to the results, the interfacial bonding strengths normal to a grain boundary are approximately 116 and 113 MPa, respectively. It is found that delamination happens in considerable low level for Jin shown in Fig. 6. The followings are concluded; transition from the plane strain to plane stress condition by parting into thin plate at a main crack-tip, and the stress relaxation with delamination the exogenous factor will become main factor on the toughness enhancement at low temperature of this alloy. In addition, the author analyzed the internal stress of inclusion particles based on Eshelby model and HRR stress singularity in the A2091 alloy. In this alloy, it was confirmed that the coarse particle of CuAl2 and Al2CuMg were fractured actually in the in-situ SEM observation at a main crack-tip, and it clarified that the fracture strength of 5–8 mm particle is about 710MPa (Fig. 7) [16]. Without limiting to Al–Li alloy, such observation is regarded as bringing about the suggestion which is useful for the development of high toughness aluminum alloys in the future.
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
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3.2. Other wrought alloys Al–Mg–Si (A6061) alloy is expanding its use for good extrudability. On this alloy, the author’s result on the effect of quenching rate after solution treatment, (width of PFZ) on fracture toughness characteristics is shown in Fig. 8. Toughness (JIC as an initiation resistance and Tmat as a propagation resistance) is lowered clearly with the lowering of the quenching rate. According to the in-situ SEM observation, it was recognized that the coarse Mg2Si particles are broken (or debonded) at several 100 mm ahead of a main crack-tip. According to the analysis, the fracture strength of Mg2Si particle of 1 – 1.51 mm diameter was about 440 MPa [17]. Stress–strain curves at the impact tensile test of this alloy are shown in Fig. 9. The strength rises by increasing strain rate; on the other hand, the elongation also shows the increase mainly by the uniform elongation [18]. By the way, there is a barrier as a practical material even in the case of 7000 system alloys which can realize
Fig. 6. Load-deflection curves at (a) liquid helium and (b) room temperatures in the fracture toughness test, illustrating the predicted onset of delamination in respective temperature. (2091 alloy)
Fig. 7. Estimated fracture strengths of CuAl2 and Al2CuMg particles as a function of diameter. (2091 alloy).
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T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
the highest strength around 600 – 650 MPa in aluminum alloys, but there is not large innovation in the strength level for a long time. The author and co-worker [19] have succeeded in a realization of high strength that exceeds 1 Gpa in Al-20% Nb heavily worked wire recently. It was produced by avoiding the generation of noxious intermetallic compound and introducing BCC metal into FCC metal by powder metallurgy route with the aim of generating the deformation mode of Nb in Al under plane strain condition. Fig. 10 shows the relationship between tensile strength and drawing strain, h. When h is 14.6, tensile strength reaches 1
Fig. 9. Stress – strain curves at three different strain rates. (6061-T6)
Gpa, and Nb shows the wavey sheet shape in SEM image. The distance between pieces is approximately 140 nm, and it seems to be a primary reason of such high strength because this effectively stops the movement of dislocations. Al–Fe or Al–Cr system may be similarly treated.
4. Casting alloys
Fig. 8. Effect of quenching rate on fracture toughness parameters of 6061-T6 alloy.
Al–Si casting alloys are widely used for engineering applications especially in automobile industry. Data on mechanical properties of such alloys, however, are not yet enough and many studies are now required. Generally, aluminum casting alloys, like other casting alloys will form dendrite during solidification. Fracture occurs generally in ductile manner by void initiation at eutectic silicon particles and inclusions in Al–Si system alloys [20]. Strength and elongation increase with decreasing the volume fraction of the second phase, especially in elongation. However, mechanical properties of the commercial Al–Si casting alloys are known to be largely dependent on the eutectic silicon characteristics (i.e. size and morphology) as well as the eutectic silicons spacing. Mechanical properties of the casting alloys are also known to be largely dependent on the dendrite arm spacing and cell size. As dendrite cell size decreases, tensile strength and elongation increase. Toughness of Al-0.5% Fe casting alloy is dependent on dendrite cell size rather than grain size, dendrite cell size becomes finer, resulting in the higher toughness [21]. However, according to the recent proposal [22], [23], mechanical properties are dominated directly by eutectic phase distribution. The structural integrity parameter, f, therefore, has been presented (Fig. 11). The relationship between the dynamic fracture toughness and the dendrite arm spacing (l2) is shown in Fig. 12(a). In Fig. 12(b), the fracture toughness data is related to the
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
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Fig. 12. Relationship between fracture toughness and secondary dendrite arm spacing and factor in high purity Al-8%Si casting alloy.
Fig. 10. Ultimate tensile strength as a function of drawing strain. (Al-20% Nb wire)
Fig. 11. Schematic illustration of the microstructural parameters considered in hypocutectic Al–Si casting alloy.
structural integrity parameter f (= MFP/l2). Better correlation between the dynamic fracture toughness as well as the other mechanical properties and the parameter has been noted [24]. Therefore the proposed structural integrity parameter is a promising factor having an influential role in controlling the mechanical properties and the deformation behavior of hypoeutectic Al– Si alloy at certain stages in the fracture process. Iron is the most commonly acquired impurity element in aluminum casting alloys affecting the fracture toughness due to the formation of needle-like Fe–Al based intermetallic compounds. The amount of needlelike Fe–Al based intermetallic compound increase with Fe content in aluminum casting alloys by recycling, and they degrade the mechanical properties and fracture toughness. The effects of iron on the dynamic fracture toughness, Jd and tearing modulus, Tmat, of the Al–Si– Cu system aluminum casting alloy are shown in Fig. 13. It is found from Fig. 13 that Jd and Tmat decrease with the increasing Fe content. However, it has been clarified that Jd and Tmat are improved when Ca is added to AC2B (Al–Si–Cu)-T6 [26]. In order to achieve good toughness, morphologies of eutectic silicon and needlelike Fe–Al or Fe–Al–Si type intermetallic compounds should be modified and heat treated [25,26]. The remarkable microstructural changes in hypereutectic Al–Si alloys can be produced by treating the melt with additives that introduce a small amount of phosphorus. In order to improve the toughness, Si particles must be refined and spheroidized. For many years, sodium (Na) has been the prevailing modifying agent
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T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
for hypocutectic Al – Si casting alloys. However, the problems associated with the use of Na as a modifier attracted many investigators to research for alternative element(s) for modifying Al – Si alloys. Recently, strontium (Sr) or stibium (Sb) has been accepted as a modifier, respectively. Jc and Jd of AC4CH(Al–Si– Mg)-T6 are strongly improved by adding Sr. These improvements in Jc and Jd can be attributed to the relative fineness and scale down of eutectic Si particles by modification [27]. The relationship between fatigue crack propagation rate, da/dN and stress intensity factor range, DK in various Al–Si casting alloys is shown in Fig. 14. Crack propagation rate is the fastest in Al-20% Si alloy in the higher DK range. According to the measurement of closure effect, Dkeff.th increases with increasing
Fig. 15. The da/dN curves obtained from impact and usual fatigue tests of AC2B-T6 casting alloys.
Fig. 13. Tmat and dynamic fracture toughness versus Fe content in AC2B-T6 casting alloy with and without Ca addition.
of Si content in as-cast and solution treated alloys [28, 29]. Many components are often operated under repeated impact loading conditions. However, such data have been reported little. The relationship between impact and usual fatigue crack propagation rate, da/dN and nominal cyclic stress intensity factor range, DK in AC2B-T6 casting alloy containing 0.2mass%Fe content is shown in Fig. 15. The impact fatigue properties are largely reduced and the retardation of da/dN is pronounced when Ca is added in this alloy [30].
5. MMCs
Fig. 14. Relationship between da/dN and DK in Al–Si casting alloys.
Metal matrix composites (MMCs) are expected as next age advanced materials. However its low toughness must be overcome. Many researchers have reported that when a main crack was propagated in the discontinuously reinforced MMCs, numerous microcracks were initiated ahead of the main crack, and these were caused by breakage of reinforcements and/or reinforcement-matrix interfacial debonding. It suggests that such crack propagation mechanisms accompanying microcracking have significant influences on the fracture of the MMCs. Such crack growth mechanisms were described in terms of the fracture mechanics and a computer simulation program for crack growth in the MMCs were constructed [31]. When microcrack exists ahead of the main crack, the crack will be deflected as shown in Fig. 16. When the main crack passes through the microcrack (dotted line),
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
anti-shielding and then shielding effects will appear. Based on such suggestions, some segregated SiC whiskers made from spray-drying method were embedded into the 6061 alloy matrix. Fig. 17 shows the
Fig. 16. Crack path morphologies when a crack passes by a single microcrack (6061-22% SiCw).
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expected good result [32]. Microstructure of the developed agglomerated MMC (MT) is also shown. Some literature reports an increase in strength to a greater or lesser degree if the clustering phase contains a sufficient number of reinforcements. This suggests the possibility of optimising the strengthening ratio of composites by controlling the spatial distribution pattern of reinforcements. However, the conditions and the mechanisms in which such increase in strengthening ratio was obtained have not been addressed. Therefore, for computational purposes, two-dimensional idealised microstructures in which one phase was a network and another isolated like islands were analysed using a plane strain formulation by FEM [33]. Fig. 18 shows the strengthening ratio of the composites as a function of the ratio of the secant Young’s moduli of the harder phase to the softer phase. Both the strengthening ratios of the composite and the secant Young’s moduli are defined at an overall strain level of 1 and 0.4% in Fig. 18(a) and (b), respectively. Fig. 18a gives the results for 25%SiCp/99.99%Al system and Fig. 18b for 25%SiCp/A2124-T6 aluminum alloy system. This suggested the existence of the optimum degree of secant Young’s modulus ratio for composite strengthening as shown by dotted lines in Fig. 18. The absolute value of the optimum degree of clustering L seems to lie around secant modulus ratio, E H S /E S =100 or above according to both Fig. 18(a) and (b). Of course, since the overall strength of the composite also depends on the volume fraction of the reinforcement, which has not been considered in the analysis, therefore quantitative interpretation as in Fig. 18(a) is not so simple and straightforward. However, it is worth noting that the fabricated agglomerated MMC provided a closer agreement with the above mentioned analysis.
6. Summary Strength and fracture problems of aluminum alloys have been stated mainly based on the author’s results. Their use will be increased in the next century. Mesomechanics considering about microstructures will become more important. Innovation in the strength level of aluminum alloys is desired; for this purpose, we must clarify fracture mechanism more in its detail.
References
Fig. 17. Representing load-displacement curves of fracture toughness test (a) and microstructure of segregated MMC (MT) (b).
[1] K.H. Schwalbe, Eng. Frac. Mech. 9 (1977) 795. [2] H.J. Kim, F. Kobayashi, M. Niinomi, T. Hagiwara, T. Sakamoto, Mat. Sci. Forum 217 – 222 (1996) 1499 –1504. [3] T. Kobayashi, M. Niinomi, MRS Int. Mtg. Adv. Mat. 5 (1989) 379 – 384.
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[4] J.R. Rice, D.M. Tracey, J. Mech. Phys. Sol. 17 (1969) 201. [5] F.A. McClintock, J. Appl. Mech. 35 (1968) 363. [6] L.M. Brown, J.D. Embury, Proceedings of the Third Internatioal Conference on the Strength of Metals and Alloys, ISIJ, London, 1975. [7] K Welpmann, A. Gysler, G. Lutjering, Z. Metallk. 71 (1980) 7. [8] G.G. Garett, J.F. Knott, Met. Trans. 9A (1978) 1187. [9] R.O. Ritchie, A.W. Thompson, Met. Trans. 16A (1985) 233. [10] G.T. Hahn, A.R. Rosenfield, Met. Trans. 6A (1975) 653. [11] T. Kobayashi, M. Niinomi, K. Degawa, Mat. Sci. Techn. 5 (1989) 1013 – 1019. [12] K.T. Venkateswara Rao, W. Yu, R.O. Ritchie, Met. Trans. 20A (1989) 485. [13] J. Glazer, S.L. Verzasconi, R.R. Sawtell, J.W. Morris Jr, Met. Trans. 18A (1987) 1695. [14] K. Welpmann, Y.T. Lee, M. Peters, Mat. Sci. Eng. A129 (1990) 21. [15] A. Takahashi, T. Kobayashi H. Toda, Int. J. Mat. Prod. Tech. 14 (1999) in press. [16] H. Toda, T. Kobayashi A. Takahashi, Alum. Trans. 1 (1999) 109 – 116. [17] S. Morita, T. Kobayashi, H. Toda, in press. [18] S. Morita, T. Kobayashi, H. Toda, Proc. ICAA6 2 (1998) 1123 – 1128. [19] T. Kobayashi, H. Toda, Proc. ICAA 7 (2000) in press.
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[20] T. Kobayashi, Proc. ICA 1 (1998) 127 – 138. [21] S. Nishi, T. Kobayashi, F. Nonoyama, J. Jpn. Inst. Met. 41 (1977) 319. [22] M.F. Hafiz, T. Kobayashi, Proceedings of the Fourth International Conference on Al Alloys, vol. 1, 1994, 107 – 114. [23] M.F. Hafiz, T. Kobayashi, Script. Met. 30 (1994) 475. [24] M.F. Hafiz, T. Kobayashi, N. FataHalla, Cast Met. 7 (1994) 103. [25] T. Kobayashi, H.J. Kim, M. Niinomi, Mat. Sci. 13 (1997) 497. [26] T. Kobayashi, M. Niinomi, M. Yamaoka, Y. Shimomura, J. Jpn Inst. Light Met. 43 (1993) 472. [27] Casting Defects and Fracture Toughness of Aluminum Casting Alloys, Japan Association of Light Metals, 1992 [28] H. Toda, T. Gouda T. Kobayashi, Mat. Sci. Techn.,14 (1998), 925 – 932. [29] T. Kobayashi, M. Niinomi, H. Egashira, T. Harata, Proceedings of the Strength of Ductile Cast Iron and other Cast Metals. Jap. Soc. Mech. Eng., Tokyo, 1993, pp. 213 – 218. [30] T. Kobayashi, H.J. Kim, M. Niinomi, Mat. Sci. Techn. 13 (1997) 497 – 502. [31] H. Toda, T. Kobayashi, Met. Mat. Trans. 28A (1997) 2149. [32] T. Kobayashi, H. Toda, Mat. Sci. Forum 242 (1997) 193– 198. [33] H. Toda, T. Gouda, T. Kobayashi, Mat. Sci. Techn. 14 (1998) 925 – 932.
Strength and fracture of aluminum alloys T. Kobayashi * Department of Production Systems Engineering, Toyohashi Uni6ersity of Technology, Tempaku, Toyohashi 4418580, Japan Received 14 June 1999
Abstract In aluminum alloys, it is general that dimple type fracture occurs from inclusions or second phases particles. Intergranular, local shear and delamination type fractures are also sometimes observed. On Al – Li alloy (A2091), effect of inclusion particles on fracture behavior is analyzed using HRR singularity and Eshelby type internal stress analysis. Fracture of inclusions of CuAl2 and Al2CuMg is observed (5–8 mm) and their strength is estimated about 710 MPa and decreased with increasing particle size. Fracture of Al3Zr or Al3Ti particles is not observed. Moreover, it is clarified that delamination occurs when the normal stress against the grain boundary attains 116 and 113 MPa at liquid He and room temperature, respectively. Moreover, newly developed high strength Al–Nb wire with 1 Gpa is introduced. Fracture and fatigue behaviors of aluminum casting alloys are also stated. Increasing the iron content, fracture toughness and fatigue properties (also impact fatigue properties) are degraded. The effect of Ca addition that possesses the modifying effect of the eutectic Si morphology is examined. It has been observed that the harmful effect of iron is largely improved. Increasing the Si content in Al – Si casting alloys, it is generally observed that fatigue strength or DKth is increased. Toughness of MMC is generally rather low. Based on the simulation analysis on fracture of Al–SiCW composite, high toughness MMC has been already developed, where some aggregated SiCW granules have been embedded into A6061 matrix. According to FEM analysis, it has been suggested that there is an appropriate aggregation ratio in the granule (i.e. SiCW content in the granule). The above topics have been reviewed. Published by Elsevier Science S.A. Keywords: Fracture toughness; Al–Li alloy; Aluminum casting alloy; Al – Nb wire, MMCs
1. Introduction The researches of strength and fracture of aluminum alloys have been actively carried out mainly in the aircraft industry related field. Generally ductile fractures occur with the sequence of nucleation, growth and coalescence of voids. It is very important to accurately understand the relationship between fracture behavior and microstructure for the purpose of alloy design. In aluminum alloys, generally, there is no ductile– brittle transition phenomenon observed in BCC metals. This provides a problem for improving the ductile fracture resistance when the yield strength is increased. Generally the toughness of the overaged alloys tends to decrease in comparison with the underaged ones. The unstable fast fracture becomes frequent, even if it is a ductile fracture, because the stengthening lowers the
* Tel.: +81-532-44-6693; fax: + 81-532-44-6690. E-mail address: [email protected] (T. Kobayashi) 0921-5093/00/$ - see front matter Published by Elsevier Science S.A. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 6 4 9 - 8
level of toughness. This becomes a problem with largescale structures. In addition, fracture characteristics under the impact load seem to become also important, because the application to automobiles and vehicles will increase. It is also predicted that the use of casting materials and aluminum matrix composites will also be expanded. It is, therefore, important to accumulate the data on the safety for the fracture behavior of aluminum alloys. In this review paper, it is reported on the problems of strength and fracture from the above viewpoints for wrought and casting aluminum alloys and composites, mainly based on the author’s studies.
2. Features of strength and fracture in aluminum alloys Since many second phase particles are contained, the fracture of aluminum alloys is the ductile fracture of the dimple formation type. In this case, whether it forms the dimple by fracture of the particle itself or by debonding at the particle interface, will change accord-
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
ing to its properties, shapes, volume fraction and condition of matrix interface of the particle. The outline of ductile fracture of void formation type is generally described in the following text. The ductile fracture progresses generally by the process of nucleation-growth-coalescence of the voids. Fig. 1 shows the various types of void in such case [1]. The nucleus of the void is inclusion, dispersoid or precipitate particle. In titanium alloys, such inclusion particles often do not exist. It is considered that in this case, void is nucleated by the mechanism of interference of grain boundary and slip band shown in Fig. 1(d) and (e). The void grows, when it is nucleated, the parameter on stress triaxiality sm /s¯ :sm is hydrostatic component of tensile stress, s is effective stress of sm) is very important. It is possible to obtain the stress triaxiality from the equation of Bridgeman in the tensile test [2], [3]. sm/s=1/3 + ln(a/2R +1)
(1)
where a is the smallest radius in void growth in the tensile test specimen, R is a projection radius of curvature. It is related to the stress triaxiality, Rice and Tracey give the following equation as Rv in respect of the void growth rate [4].
Fig. 1. Schematic illustration of void nucleation. Schwalbe et al. [1].
dRv/Rv = 0.28 dop exp(1.5sm/s)
(2)
where op is an effective plastic strain. Fig. 2 shows the results of examining this relation in 7075 system aluminum alloys [3]. It is clear that the ratio of void growth increases by increasing the stress triaxiality, and there is the rectilinear relation. It is also proven that coefficient of 0.28 differs by the alloys. It is shown that the void is easy to grow regardless of the existence of nucleus of the void, when the stress triaxiality increases. The coalescence of voids has not always been clarified, since it is unstably and rapidly occurred. McClintock [5] considered that adjoining voids grow and coalesce. On the other hand, the small voids may support coalescence shown in Fig. 1(b). From a past, however, the mechanism coupled by the shear by the local strain concentration (formation of void sheet) between voids has been indicated. Fig. 1(f) indicates that the straight line, which connects the void root, becomes p/4 rad (45°), when the void length (depth) becomes equal to a distance between voids [6]. Fig. 3 shows deformation models in the age-hardenable alloys [7]. In age-hardenable aluminum alloys, precipitate particles (0.02–0.1 mm) are sheared by the dislocations. By inducing the coarse slip over the grain size, this may cause the embrittlement (type A). In type B, by precipitating the particle and forming PFZ (precipitate free zone) in grain boundary vicinity, and coarse particles precipitate on grain boundary, the intergranular fracture is easy to be caused. In Al–Li alloy, the fracture patterns of A and B types are mainly observed. Especially, there is the high possibility of connecting with the macroscopic shear fracture in case of A type. However, when to some extent large dispersoid particles exist like type C, toughness is improved. In case of B type, the local stress concentration at PFZ increases when the difference of (sM –sPFZ) increases, where the strengths of matrix and PFZ define sM and sPFZ, respectively, the fracture tendency at PFZ increases. However, most large factor is void nucleation at coarse particle and after extension growth on grain boundary. For the fracture of void formation type like the above, Garrett et al. [8] give the following equation as a result of occurring the fracture, when the strain at a main crack-tip becomes threshold value, o*. c 2 2 1/2 KIC : (2CEo*·s c yn /1−6 )
Fig. 2. Relationship between void growth rate and stress triaxiality in 7075 aluminum alloys. (Zr content (mass%) AB 0.01, B:0.06, C:0.08, D:0.16, E:030)
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(3)
where C is constant (= 1/40), sy is yield stress, n is work hardening exponent and 6 is Poisson ratio. On the other hand, Ritchie et al. [9] have indicated that the strain at a main crack-tip must have exceeded limitative fracture strain o*(s c m/s) over a characteristic distance l*0 (order of process zone= CTOD) in case of such strain dominant fracture. In 3-point bending test on above-mentioned 7075 system aluminum alloy (E),
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3. Wrought alloys
3.1. Al–Li alloys
Fig. 3. Deformation and fracture model of age-hardenable alloys. (Welpmann et al. [7]).
In Al–Li alloy, the age hardening occurs by the precipitate process below. SSSS G.P. zone d%(Al3Li)d(AlLi)
Fig. 4. Distribution of plastic strain and critical fracture strain as a function of stress triaxiality at a main crack-tip. X is the distance from a main crack tip, x1 is the first void nucleation point (7075 alloy).
Fig. 5. Variation of maximum load (Pm) and absorbed energy (Et) as a function of test temperature obtained by CAI system in 8090 aluminum alloy.
the example in search of ratio between distance from a main crack-tip, x and CTODIC (dIC) and relationship between local effective plastic strain, op is shown in Fig. 4. It has been confirmed that intersection point, xI between op and o*(s f m/s) corresponds to distance, l between second phase particles in the first void nucleation point [3]. The void width and depth are defined as W and h, and the roughness of fracture surface is defined as M =h/M, of = ln(M 2/3f)/3 and, JIC =s0ofl*0 can be considered ( f: volume fraction of particle,s0: flow stress), JIC :s0/3 ln(M 2/3f)·l*. 0
(4)
On the other hand, Hahn et al. [10] give following equation, the crack occurs, when the second phase particle comes in the process zone. KIC :[2syE(p/6)1/3d]1/2f − 1/6
(5)
where E is Young’s modulus, d is diameter of second phase particle. Though this equation agrees with the experimental value, the contradiction has also been indicated for KIC rising in proportion to d 1/2 and s 1/2 y .
The globular and coherent d% phase mainly contributes to the hardening. However, d% phase is easy to be cut off by dislocations, and the sedimentation of the dislocation to the grain boundary is easy to pile up. The local shearing deformation is easily frequent, and there is the high possibility of separation and intergranular fracture in the slip plane. In addition, in this alloy, the absorption of the hydrogen gas is frequent, and low toughness is a problem. However, this alloy shows the positive temperature dependence of the strength observed in Ni based superalloys, and it is excellent for the high temperature strength. Whereas, it is excellent in toughness under low temperature, and it has also been applied to the fuel tank of space shuttle. Fig. 5 shows the example on such enhancement of low-temperature toughness [11]. This is related to the delamination which originates from the fibrous structure by the rolling, and the effect of the anisotropy seems to be big [12]. Ritchie et al. [12] have asserted that such exogenous factor is a primary cause on the enhancement of low-temperature toughness [13,14]. Recently the author and co-workers have examined the initiation stress of delamination crack of the A2091 alloy (L-T direction) by the ultrasonic wave measurement at room temperature and liquid helium temperature [15]. According to the results, the interfacial bonding strengths normal to a grain boundary are approximately 116 and 113 MPa, respectively. It is found that delamination happens in considerable low level for Jin shown in Fig. 6. The followings are concluded; transition from the plane strain to plane stress condition by parting into thin plate at a main crack-tip, and the stress relaxation with delamination the exogenous factor will become main factor on the toughness enhancement at low temperature of this alloy. In addition, the author analyzed the internal stress of inclusion particles based on Eshelby model and HRR stress singularity in the A2091 alloy. In this alloy, it was confirmed that the coarse particle of CuAl2 and Al2CuMg were fractured actually in the in-situ SEM observation at a main crack-tip, and it clarified that the fracture strength of 5–8 mm particle is about 710MPa (Fig. 7) [16]. Without limiting to Al–Li alloy, such observation is regarded as bringing about the suggestion which is useful for the development of high toughness aluminum alloys in the future.
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
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3.2. Other wrought alloys Al–Mg–Si (A6061) alloy is expanding its use for good extrudability. On this alloy, the author’s result on the effect of quenching rate after solution treatment, (width of PFZ) on fracture toughness characteristics is shown in Fig. 8. Toughness (JIC as an initiation resistance and Tmat as a propagation resistance) is lowered clearly with the lowering of the quenching rate. According to the in-situ SEM observation, it was recognized that the coarse Mg2Si particles are broken (or debonded) at several 100 mm ahead of a main crack-tip. According to the analysis, the fracture strength of Mg2Si particle of 1 – 1.51 mm diameter was about 440 MPa [17]. Stress–strain curves at the impact tensile test of this alloy are shown in Fig. 9. The strength rises by increasing strain rate; on the other hand, the elongation also shows the increase mainly by the uniform elongation [18]. By the way, there is a barrier as a practical material even in the case of 7000 system alloys which can realize
Fig. 6. Load-deflection curves at (a) liquid helium and (b) room temperatures in the fracture toughness test, illustrating the predicted onset of delamination in respective temperature. (2091 alloy)
Fig. 7. Estimated fracture strengths of CuAl2 and Al2CuMg particles as a function of diameter. (2091 alloy).
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the highest strength around 600 – 650 MPa in aluminum alloys, but there is not large innovation in the strength level for a long time. The author and co-worker [19] have succeeded in a realization of high strength that exceeds 1 Gpa in Al-20% Nb heavily worked wire recently. It was produced by avoiding the generation of noxious intermetallic compound and introducing BCC metal into FCC metal by powder metallurgy route with the aim of generating the deformation mode of Nb in Al under plane strain condition. Fig. 10 shows the relationship between tensile strength and drawing strain, h. When h is 14.6, tensile strength reaches 1
Fig. 9. Stress – strain curves at three different strain rates. (6061-T6)
Gpa, and Nb shows the wavey sheet shape in SEM image. The distance between pieces is approximately 140 nm, and it seems to be a primary reason of such high strength because this effectively stops the movement of dislocations. Al–Fe or Al–Cr system may be similarly treated.
4. Casting alloys
Fig. 8. Effect of quenching rate on fracture toughness parameters of 6061-T6 alloy.
Al–Si casting alloys are widely used for engineering applications especially in automobile industry. Data on mechanical properties of such alloys, however, are not yet enough and many studies are now required. Generally, aluminum casting alloys, like other casting alloys will form dendrite during solidification. Fracture occurs generally in ductile manner by void initiation at eutectic silicon particles and inclusions in Al–Si system alloys [20]. Strength and elongation increase with decreasing the volume fraction of the second phase, especially in elongation. However, mechanical properties of the commercial Al–Si casting alloys are known to be largely dependent on the eutectic silicon characteristics (i.e. size and morphology) as well as the eutectic silicons spacing. Mechanical properties of the casting alloys are also known to be largely dependent on the dendrite arm spacing and cell size. As dendrite cell size decreases, tensile strength and elongation increase. Toughness of Al-0.5% Fe casting alloy is dependent on dendrite cell size rather than grain size, dendrite cell size becomes finer, resulting in the higher toughness [21]. However, according to the recent proposal [22], [23], mechanical properties are dominated directly by eutectic phase distribution. The structural integrity parameter, f, therefore, has been presented (Fig. 11). The relationship between the dynamic fracture toughness and the dendrite arm spacing (l2) is shown in Fig. 12(a). In Fig. 12(b), the fracture toughness data is related to the
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
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Fig. 12. Relationship between fracture toughness and secondary dendrite arm spacing and factor in high purity Al-8%Si casting alloy.
Fig. 10. Ultimate tensile strength as a function of drawing strain. (Al-20% Nb wire)
Fig. 11. Schematic illustration of the microstructural parameters considered in hypocutectic Al–Si casting alloy.
structural integrity parameter f (= MFP/l2). Better correlation between the dynamic fracture toughness as well as the other mechanical properties and the parameter has been noted [24]. Therefore the proposed structural integrity parameter is a promising factor having an influential role in controlling the mechanical properties and the deformation behavior of hypoeutectic Al– Si alloy at certain stages in the fracture process. Iron is the most commonly acquired impurity element in aluminum casting alloys affecting the fracture toughness due to the formation of needle-like Fe–Al based intermetallic compounds. The amount of needlelike Fe–Al based intermetallic compound increase with Fe content in aluminum casting alloys by recycling, and they degrade the mechanical properties and fracture toughness. The effects of iron on the dynamic fracture toughness, Jd and tearing modulus, Tmat, of the Al–Si– Cu system aluminum casting alloy are shown in Fig. 13. It is found from Fig. 13 that Jd and Tmat decrease with the increasing Fe content. However, it has been clarified that Jd and Tmat are improved when Ca is added to AC2B (Al–Si–Cu)-T6 [26]. In order to achieve good toughness, morphologies of eutectic silicon and needlelike Fe–Al or Fe–Al–Si type intermetallic compounds should be modified and heat treated [25,26]. The remarkable microstructural changes in hypereutectic Al–Si alloys can be produced by treating the melt with additives that introduce a small amount of phosphorus. In order to improve the toughness, Si particles must be refined and spheroidized. For many years, sodium (Na) has been the prevailing modifying agent
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for hypocutectic Al – Si casting alloys. However, the problems associated with the use of Na as a modifier attracted many investigators to research for alternative element(s) for modifying Al – Si alloys. Recently, strontium (Sr) or stibium (Sb) has been accepted as a modifier, respectively. Jc and Jd of AC4CH(Al–Si– Mg)-T6 are strongly improved by adding Sr. These improvements in Jc and Jd can be attributed to the relative fineness and scale down of eutectic Si particles by modification [27]. The relationship between fatigue crack propagation rate, da/dN and stress intensity factor range, DK in various Al–Si casting alloys is shown in Fig. 14. Crack propagation rate is the fastest in Al-20% Si alloy in the higher DK range. According to the measurement of closure effect, Dkeff.th increases with increasing
Fig. 15. The da/dN curves obtained from impact and usual fatigue tests of AC2B-T6 casting alloys.
Fig. 13. Tmat and dynamic fracture toughness versus Fe content in AC2B-T6 casting alloy with and without Ca addition.
of Si content in as-cast and solution treated alloys [28, 29]. Many components are often operated under repeated impact loading conditions. However, such data have been reported little. The relationship between impact and usual fatigue crack propagation rate, da/dN and nominal cyclic stress intensity factor range, DK in AC2B-T6 casting alloy containing 0.2mass%Fe content is shown in Fig. 15. The impact fatigue properties are largely reduced and the retardation of da/dN is pronounced when Ca is added in this alloy [30].
5. MMCs
Fig. 14. Relationship between da/dN and DK in Al–Si casting alloys.
Metal matrix composites (MMCs) are expected as next age advanced materials. However its low toughness must be overcome. Many researchers have reported that when a main crack was propagated in the discontinuously reinforced MMCs, numerous microcracks were initiated ahead of the main crack, and these were caused by breakage of reinforcements and/or reinforcement-matrix interfacial debonding. It suggests that such crack propagation mechanisms accompanying microcracking have significant influences on the fracture of the MMCs. Such crack growth mechanisms were described in terms of the fracture mechanics and a computer simulation program for crack growth in the MMCs were constructed [31]. When microcrack exists ahead of the main crack, the crack will be deflected as shown in Fig. 16. When the main crack passes through the microcrack (dotted line),
T. Kobayashi / Materials Science and Engineering A280 (2000) 8–16
anti-shielding and then shielding effects will appear. Based on such suggestions, some segregated SiC whiskers made from spray-drying method were embedded into the 6061 alloy matrix. Fig. 17 shows the
Fig. 16. Crack path morphologies when a crack passes by a single microcrack (6061-22% SiCw).
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expected good result [32]. Microstructure of the developed agglomerated MMC (MT) is also shown. Some literature reports an increase in strength to a greater or lesser degree if the clustering phase contains a sufficient number of reinforcements. This suggests the possibility of optimising the strengthening ratio of composites by controlling the spatial distribution pattern of reinforcements. However, the conditions and the mechanisms in which such increase in strengthening ratio was obtained have not been addressed. Therefore, for computational purposes, two-dimensional idealised microstructures in which one phase was a network and another isolated like islands were analysed using a plane strain formulation by FEM [33]. Fig. 18 shows the strengthening ratio of the composites as a function of the ratio of the secant Young’s moduli of the harder phase to the softer phase. Both the strengthening ratios of the composite and the secant Young’s moduli are defined at an overall strain level of 1 and 0.4% in Fig. 18(a) and (b), respectively. Fig. 18a gives the results for 25%SiCp/99.99%Al system and Fig. 18b for 25%SiCp/A2124-T6 aluminum alloy system. This suggested the existence of the optimum degree of secant Young’s modulus ratio for composite strengthening as shown by dotted lines in Fig. 18. The absolute value of the optimum degree of clustering L seems to lie around secant modulus ratio, E H S /E S =100 or above according to both Fig. 18(a) and (b). Of course, since the overall strength of the composite also depends on the volume fraction of the reinforcement, which has not been considered in the analysis, therefore quantitative interpretation as in Fig. 18(a) is not so simple and straightforward. However, it is worth noting that the fabricated agglomerated MMC provided a closer agreement with the above mentioned analysis.
6. Summary Strength and fracture problems of aluminum alloys have been stated mainly based on the author’s results. Their use will be increased in the next century. Mesomechanics considering about microstructures will become more important. Innovation in the strength level of aluminum alloys is desired; for this purpose, we must clarify fracture mechanism more in its detail.
References
Fig. 17. Representing load-displacement curves of fracture toughness test (a) and microstructure of segregated MMC (MT) (b).
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