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M Al–Zn–Mg alloy
Int. J. Fatigue Vol. 20, No. 10, pp. 757–766, 1998 1998. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain 0142–1123/98/$—see front matter
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Evaluation of stress corrosion resistance and corrosion fatigue fracture behavior of ultrahigh-strength P/M Al–Zn–Mg alloy Kohji Minoshima*‡, Makoto Okada† and Kenjiro Komai* *Department of Mechanical Engineering, Graduate School of Engineering, Kyoto University, Yoshida-Honmachi, Sakyo-Ku, Kyoto 606-8501, Japan †Central Japan Railway Co., Yaesu 1-6-6, Chuo-Ku, Tokyo 103-8288, Japan (Accepted 21 May 1998) Quasi-static tensile tests in air and slow strain rate tests (SSRTs) in a 3.5% NaCl solution were conducted in an ultra-high-strength P/M Al–Zn–Mg alloy fabricated through powder metallurgy. Attention is also paid to fatigue strength and fatigue crack growth behavior in laboratory air and in a 3.5% NaCl solution. The alloy has extremely high strength of about 800 MPa. However, elongation at break remains small, at about 1.3%. The final fracture occurs by a macroscopically flat crack normal to the tensile axis, with little reduction in area and little shear lip on the periphery of a smooth sample. However, it fails microscopically in a ductile manner, with dimples. Dimple size is less than 1 m, because the grain size of the alloy is extremely small. Strengthening mechanisms operating in the alloy are: small grains, sufficient metastable ⬘ phase in a matrix, and intermetallic compound acting as a fiber reinforcement. The SSRT strength in a 3.5% NaCl solution decreases slightly at a very low strain rate, that is smaller than those observed in aluminum alloys sensitive to stress corrosion. This means that the crack initiation resistance to stress corrosion is superior. However, under cyclic loading, the corrosion fatigue strength becomes lower than that conducted in air, because pitting corrosion on a sample surface acts as a stress concentrator. Crack initiation site of quasi-static and fatigue failure of the alloy is at inclusions, and hence, it is essential to decrease inclusions in the alloy for the improvement of the mechanical properties. Fatigue crack resistance of the alloy is inferior to conventional Al–Zn–Mg alloys fabricated by ingot metallurgy, because the fatigue fracture toughness, or ductility, of the alloy is inferior to other Al alloys, and intergranular cracking promotes crack growth. However, no influence of 3.5% NaCl solution on corrosion fatigue crack growth is observed, although an investigation is required into whether stress corrosion crack growth occurs or not, and at the same time, and of corrosion fatigue crack growth behavior at lower stress intensity. The fracture surface and crack initiation sites are closely examined using a high-resolution field emission type scanning electron microscope, and the fracture mechanisms of the alloy are discussed. 1998. Published by Elsevier Science Ltd. All rights reserved (Keywords: corrosion fatigue; stress corrosion cracking; SSRT; fatigue strength; fatigue crack propagation; fractography; powder metallurgy; Al–Zn–Mg alloy)
INTRODUCTION
and distribution of metastable precipitates acting as pinning centers for preventing dislocation motion. In particular, Mn and a very small amount of Ag additions greatly improve the tensile strength of the P/M alloys; the tensile strength exceeding 900 MPa3 is reported. However, little is known about its failure mechanism. Besides, in order to facilitate the application of these materials to machines and structures, the mechanical properties, including fatigue in various environments, have to be investigated4,5; the resistance of the materials to environments is extremely important. Important issues about mechanical properties in service are: quasi-static strength, ductility or fracture toughness, stress corrosion cracking (SCC) resistance, and also fatigue strength. The crack initiation resistance
Due to further progress in science and technology, new high-strength materials have become especially important; among these are metallic materials manufactured by powder metallurgy (P/M). One of the main attractions of P/M alloys is the ability to fabricate complex parts in an economical way: the ability of homogeneous microstructures and of element design of an alloy which is impossible by traditional ingot metallurgy (I/M). Recently, remarkable improvement in mechanical properties of P/M Al–Zn–Mg–Cu alloys has been achieved1–3 by controlling the morphology ‡Corresponding author.
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of stress corrosion cracking in a P/M Al alloy is reported to be superior to I/M aluminum alloy6,7, because a P/M alloy is subjected to rather general corrosion not to localized corrosion7. One reason is that large precipitates including Fe or Cu are not present in a P/M alloy, which act as cathodic sites to promote localized corrosion. However, the crack growth rate of a P/M alloy is almost the same as a traditional I/M Al alloy7,8. When these alloys are to be used in service, the properties in an operating environment under dynamic loading such as corrosion fatigue must be clarified. In this investigation, quasi-static tensile tests, slow strain rate tests (SSRTs) in a 3.5% NaCl solution, and fatigue strength and fatigue crack growth behavior in laboratory air and in a 3.5% NaCl solution were investigated in an ultra-high-strength P/M Al–Zn–Mg–Cu– Mn–Ag alloy3. Fracture surface and specimen surface after testing were closely examined using a field-emission type high-resolution scanning electron microscope, and the fracture mechanisms are discussed. EXPERIMENTAL PROCEDURES The material tested was a high-strength Al–9Zn–3Mg– 1.5Cu–4Mn–0.5Zr–0.04Ag alloy, fabricated by powder metallurgy3. The powders prepared by air atomizing were pressed into a rod-like shape by a cold isostatic pressing. The rod was first preheated in nitrogen gas at 500°C for 1 h, and extruded to an extrusion ratio of 20. The material was then machined to a specimen and heat-treated to a T6 condition: the specimens were solution-treated at 490°C for 2 h, then water-quenched, and aged at 120°C for 24 h. The shape and dimensions of test samples are shown in Figure 1: smooth round specimens shown in Figure 1(a) were machined in a LL orientation, where the loading axis paralleled the extrusion direction. To investigate crack growth behavior, compact tension
type specimens shown in Figure 1(b) were machined in T–L and L–T orientations. Four lots were used to prepare the specimens: see Table 1. Round smooth specimens (Lots 1, 2 and 3) were machined from a rod 40 mm in diameter, whereas CT specimens were machined from a square rod of 40 mm each side. A smooth sample surface was polished with #1500 emery paper followed by final finish with diamond paste of 2.5 m. Tensile tests were conducted in laboratory air at a displacement rate of 2 mm/min by using a computer-controlled tensile testing machine (Shimadzu Co. Ltd., Autograph AG-10TD, load capacity: 100 kN). To evaluate the SCC susceptibility, slow strain rate tests (SSRTs) were conducted at a displacement rate of 0.05 and 0.005 mm/min. A corrosive environment selected was a 3.5% NaCl solution prepared with reagent grade NaCl and deionized water (specific resistance > 1 M⍀ cm). The temperature of the solution was kept at 25°C by a thermoregulator, and was circulated by a vane pump made of synthetic resin between a corrosion chamber and a corrosion reservoir. In a 3.5% NaCl solution, a selected middle part of a sample of 2 mm in length was exposed to the environment, and the rest was coated to prevent the specimen surface from corrosion. Axial tension–tension fatigue tests at a stress ratio R of 0.1 were conducted on a smooth specimen in air and in a 3.5% NaCl solution at 25°C. In fatigue tests in air, a sinusoidal waveform at a stress cycle frequency f of 1–10 Hz was used, whereas in a 3.5% NaCl solution, a triangular wave at f = 0.1 Hz was selected. Similarly to in the SSRTs, a middle part of 2 mm in length was exposed to the environment. The testing machine employed was an electro-hydraulic fatigue testing machine with load capacity of 100 kN. Fatigue crack growth characteristics at a stress ratio R of 0.1 were evaluated by using a compact tension type specimen shown in Figure 1(b). In laboratory air, the effect of crack plane orientation on fatigue crack growth was investigated using T–L and L–T specimens. For L–T specimens, the influence of 3.5% NaCl solution was studied. In air, a sinusoidal waveform at f = 1–5 Hz was used, whereas a triangular waveform at f = 0.1 Hz was selected in a 3.5% NaCl solution. The testing machine employed was a computer-controlled electro-hydraulic fatigue testing machine (load capacity: 50 kN). Crack opening stress intensity was measured by an elastic unloading compliance technique: a strain gage was adhered to the back surface of a CT specimen, and a crack opening stress intensity was computed by using a personal computer linked with the fatigue testing machine. A surface crack length was monitored by using a traveling microscope. To compensate the tunneling effect, where a crack grew more in the inside
Table 1 P/M aluminum alloy for each set of experiments. Four lots were used to prepare the samples Smooth round specimens
Figure 1 Shape and dimensions of test specimens. All dimensions are in mm. (a) Smooth round specimen; (b) compact tension type specimen
CT specimens
Tensile tests in air
Lots 1, 2
SSRTs Fatigue tests in air Corrosion fatigue tests
Lot 1 Lots 1, 2, 3 Lots 1, 2 Lot 4
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of a specimen than on the surface, an average crack length was adopted to calculate a crack growth rate as well as stress intensity factor. An average crack length was defined by crack area divided by the sample thickness. EXPERIMENTAL RESULTS AND DISCUSSIONS Quasi-static tensile fracture behavior in air Table 2 summarizes the mechanical properties in laboratory air. The tensile strength went up to about 800 MPa which was slightly lower than the value reported in this material3. However, the strength was extremely high compared with ingot-metallurgy (I/M) 7XXX series aluminum alloys (B = 570–595 MPa: 7075-T6, 7175-T66)9. We must note that the elongation at the break was only 1.3%, which was smaller than I/M 7XXX series aluminum alloys (about 11% for 7075-T6 and 7175-T6)9. The fracture occurred by a macroscopic crack normal to the loading direction, and little shear lip could be seen (see Figure 2(a)): this is consistent with the alloy having small elongation at the break. The crack was initiated at an inside defect, a void or an inclusion (Figure 2(b)), whose diameter ranged from 70 to 100 m. This means that if the size of the defect could be made smaller, the mechanical properties might be increased. As shown in Figure 2(a), the alloy fractured in a brittle manner from the macroscopic view. However, the fracture surface was microscopically dominated by ductile fracture characterized by dimples (see Figure 2(c)). The strengthening mechanism of the alloy3 has been reported to be (1) small grain, less than 1 m, (2) metastable ⬘ phase in a matrix, and (3) intermetallic compound (Al20(Cu,Zn)2Mn3) that acts as a fiber reinforcement. The size of dimples observed on the fracture surface was about 0.5–1 m, and this suggested the strengthening mechanism due to small grain was achieved well. The other important strengthening mechanism that could be obtained from the fracture surface is intermetallic compound, which is shown by arrows in Figure 2(c). The intermetallic is about a few micrometers in length and the width was about a few hundred nanometers. The longitudinal direction of the intermetallic was parallel to the extruding direction, and they contributed to fiber reinforcement for the present P/M aluminum alloy. Slow-strain rate tests Figure 3 illustrates the fracture strength of SSRTs as a function of displacement rate in a 3.5% NaCl solution. The quasi-static fracture strengths in laboratory air are also plotted in the figure. Unlike the tensile test results in air, a crack that caused final failure in Table 2
Mechanical properties of the P/M aluminum alloy
Tensile strength (MPa) 797
Elongation (%) 1.31
Tensile tests were conducted in Lots 1 and 2, and two samples were used for each lot. The values listed in the table are the average values.
Figure 2 Fracture behavior of a quasi-static tensile test. (a) Macroscopic fracture morphology; (b) example of crack initiation site (inside inclusion); (c) fracture surface imaged with SEM. Arrows show examples of an intermetallic compound (Al20(Cu,Zn)2Mn3) that contributes to fiber reinforcement
a 3.5% NaCl solution was initiated at the sample surface. Figure 4 shows SEM photographs of a crack initiation site: a crack was initiated at a pitting corrosion site (Figure 4(a)), irrespective of loading condition, which will be refereed to as ‘pitting corrosion’. In some cases, severe corrosion occurred, and the size of the corrosion defect became still larger than pitting corrosion: this severe corrosion is presented in Figure 4(b). In Figure 3, symbols with a superscript * indicate the crack initiation site was such a large corrosion defect. When the crack initiation site was pitting corrosion, the strength in a 3.5% NaCl solution at 0.05 mm/min, which corresponded to 1.5–2.0 × 10−6 l/s,
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Figure 3 Fracture strength as a function of displacement rate in a 3.5% NaCl solution and in air. A superscript * indicates the crack initiation site was a large corrosion defect shown in Figure 4(b)
Figure 4 Crack initiation site of an SSRT in a 3.5% NaCl solution. (a) Pitting corrosion type (displacement rate: 0.005 mm/min, B = 732 MPa); (b) large corrosion defect type (displacement rate: 0.05 mm/min, B = 587 MPa)
was the same as that in laboratory air. Even in this case, the crack was initiated at a corrosion pit, not at an inside defect. If more tests were conducted at 0.05 mm/min, some inside defect could cause the crack initiation. However, at 0.005 mm/min corresponding to 1.5–2.0 × 10−7 l/s, the strength decreased from air data. In a material/environment system sensitive to stress
corrosion cracking, SSRT strength decreased with decreasing loading rate, and a minimum strength exists at a certain loading rate. Such a loading rate was reported to be 10−5 to 10−6 l/s10: this is larger than the loading rate of 0.005 mm/min (1.5–2 × 10−7 l/s) where the SSRT strength of the alloy decreased from air data. SSRTs evaluate the sensitivity of SCC in a given environment, in particular, the strength of crack initiation and passive films formed on a surface. When the crack initiation site was pitting corrosion, the strength slightly decreased at a lower loading rate of 0.005 mm/min. This decrease is small compared with that for materials sensitive to stress corrosion cracking: a decrease sometimes reaches almost half or smaller than half of the tensile strength10. This indicates that the alloy had a superior SCC strength to I/M aluminum alloys. When an aluminum alloy has very small grains such as this P/M alloy, large precipitates in a grain do not exist and, therefore, less damage of passive films is achieved, which in turn causes an increase in SCC strength. Fracture surface was dominated by dimples and corrosion products. However, at a loading rate of 0.005 mm/min, intergranular failure could be observed. The intergranular failure was typically observed in SCC of an aluminum alloy. Further investigation including crack growth of stress corrosion under a sustained load is required. When the crack initiation site was a large corrosion defect, the SSRT strength decreased from that conducted in air at both 0.05 and 0.005 mm/min, because a large corrosion defect increased the stress concentration, thereby decreasing the SSRT strength. What had caused such a severe corrosion, or corrosion defects, may be related to some kind of localized segregation due to insufficient control of the manufacturing process. Further investigation is necessary to avoid such severe corrosion: this will improve the properties of the alloy in a corrosive environment. Fatigue strength of smooth round specimens Figure 5 shows S–N curves under uniaxial tension– tension fatigue loading. The symbols with F indicate that a foreign phase could be seen in the fracture surface, which will be discussed later in detail. The fatigue results in air show a somewhat large scatter, and the specimens with foreign phases tended to have a lower fatigue strength than others. The macroscopic fracture morphology was almost the same as that fractured under quasi-static tensile loading in air: the specimen fractured with little shear lip on the periphery of the sample, and the fracture surface was normal to the axial loading direction. Under fatigue loading, a crack initiated at an inside inclusion near the specimen surface, irrespective of applied stress. Such a crack may be initiated by interfacial debonding between matrix and inclusion and/or fracture of the inclusion itself. The fracture surface consisted of fatigue and final catastrophic failure areas. The stress intensity factor obtained by using the crack length at final catastrophic failure was 7 to 8 MPa m1/2, that agreed with the fatigue fracture toughness of L– T orientation obtained by fatigue crack growth experiments, which will be shown later. Irrespective of applied stress, a fatigue crack was initiated at an inclusion near the specimen surface.
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Figure 5 S–N curves under uniaxial tension–tension fatigue loading. A superscript * indicates that the initiation site was an inclusion, while a superscript F indicates that a foreign phase could be seen in the fatigued fracture surface
Electron probe microanalysis showed that inclusions could be classified into three types: oxide inclusion, inclusion with foreign metals (Fe, Zn), and silica (SiO2). Inclusions with foreign metals or silica are considered to be brought about during manufacturing. Oxide inclusions are the most important and inherent problems for P/M metallurgy that we have to overcome. These results show that the amount of inclusions in the material should be decreased to improve fatigue strength of the present alloy. Fatigue fracture surface was similar to those observed in CT specimens shown in Figure 15. Near the crack initiation site, the fracture surface was transgranular, and macroscopically it looked flat, because the grain size of the material was extremely small. With an extension of fatigue crack, not only transgranular cracking but also intergranular cracking appeared. Fracture of intermetallic compound (Al20(Cu,Zn)2Mn3) could also be seen. At the final catastrophic failure stage, dimples as well as fracture of intermetallic compounds dominated over the fracture surface. Corrosion fatigue strength The corrosion fatigue results in a 3.5% NaCl solution are shown in Figure 5. A crack that caused final failure was initiated at either an inclusion or a corrosion pit shown in Figure 6(a). A superscript * is added to the symbol when the initiation site was an inclusion. When the corrosion fatigue life was extremely short at 300 MPa (Nf = 5.83 × 102), an extremely large corrosion-induced defect of 500 to 600 m was an initiation site: see Figure 6(b). This large corrosion defect caused a large decrease in corrosion fatigue life compared to other specimens. Except for this, the fatigue strength in a 3.5% NaCl solution was even smaller than that in air. As is discussed later, little influence of the corrosion environment on fatigue crack growth in the alloy was observed. Therefore, a decrease in corrosion fatigue
Figure 6 Crack initiation site of a corrosion fatigue test in a 3.5% NaCl solution. (R = 0.1, f = 0.1 Hz). (a) Pitting corrosion (max = 300 MPa, Nf = 1.51 × 104); (b) large corrosion defect (max = 300 MPa, Nf = 5.83 × 102)
life was due to promotion of crack initiation by pitting corrosion. Corrosion pits tended to occur along the extruding direction, and these pits coalesced with each other, resulting in a crack initiation site. In some cases, a large corrosion defect was formed on a specimen surface, shown in Figure 4(b) and Figure 6(b), resulting in a further decrease in SSRT strength or corrosion fatigue life. A reason why such large corrosion defects occurred may be due to some localized segregation of constituents: further investigation is required to improve the corrosion property which in turn improves the mechanical properties of the P/M alloy in a corrosive environment. Fatigue crack growth in air Figure 7 shows a macroscopic fatigue fracture surface of a CT specimen conducted in air: three different types of fracture morphology could be seen. These will be refered to as Type 1, Type 2, and Type 3, respectively. Respective microscopic fracture surfaces are shown in Figure 8, and the characteristic features of each fracture surface are summarized as follows: Type 1 fracture surface looked dark, and the fracture surface was completely different from the other types. In the case of L–T specimens, a crack tended to be retarded when the crack reached the Type 1 fracture surface. Type 2 fracture surface involves some typical fracture surface morphology which looked as if the mating surfaces had contacted each other. As for Type 3 fracture surface shown in Figure 8(c) at higher stress range, characteristic fracture morphology of this P/M alloy, small grains and intermetallic compound (Al20(Cu,Zn)2Mn3) acting as a reinforcement, could be seen. Type 1 and Type 2 fracture surfaces were in
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Figure 7 Macroscopic fatigued fracture surface of T–L orientation (CT specimen)
particular observed in Lot 4, that was machined into CT specimens. The different fracture morphology was due to the existence of foreign phases in the material which were elongated in the extrusion direction: the foreign phases are considered to be caused by insufficient control of a manufacturing parameter during extrusion. A ratio of each fracture surface morphology, or foreign phases, depended on a specimen and, therefore, crack growth rate varied from sample to sample. Hence, in this experiment, crack growth rates obtained by a CT specimen, of which a Type 3 fracture surface dominated over the fracture surface, are plotted. Figure 9 illustrates the fatigue crack growth rate of T–L crack plane orientation in air as a function of ⌬K. The crack growth rates of an I/M Al–Zn–Mg alloy, 7075-T611,12 and P/M Al–Zn–Mg alloy, 7090-T611 are also plotted for comparison. At lower stress intensity, there was little difference in fatigue crack growth rate. At an intermediate ⌬K, crack growth rate of T–L2 was accelerated. However, at higher stress intensity, acceleration of crack growth was observed in both T–L1 and T– L2 specimens. The fatigue fracture toughness of the alloy was about 6–7 MPa m1/2, and the crack growth rate increased as a maximum stress intensity reached the fatigue fracture toughness. We must note that crack growth rate at higher stress intensity of the present alloy was the highest, and decreased in the order of 7090-T6 and 7075-T6. This means that the fatigue fracture toughness of P/M Al alloys, in particular the present Al alloy, was inferior to conventional I/M Al alloys. Figure 10 shows the crack opening stress intensity factor as a function of maximum stress intensity factor.
Figure 8 Microscopic fatigued fracture surface of a CT specimen conducted in air (T–L orientation); (a) Type 1; (b) Type 2; (c) Type 3
The crack opening stress intensity was higher in the T–L1 specimen than the T–L2 specimen. The ratio of foreign phases of T–L1 was higher than that of the T–L2 specimen, and the roughness of fracture surface was higher in T–L1 specimen. This indicates that the crack opening stress intensity increased by roughness induced crack closure. If the ratio of foreign phases of T–L1 specimens was smaller, the crack opening stress intensity factor would be decreased.
Stress corrosion resistance and corrosion fatigue fracture behavior of P/M Al–Zn–Mg alloy
Figure 9 Crack growth rate, da/dN, as a function of ⌬K for T–L crack plane orientation
Figure 10 Crack opening stress intensity factor as a function of maximum stress intensity factor
The crack growth rates as a function of effective stress intensity factor, ⌬Keff, are shown in Figure 11. In the figure, crack growth rates of 7090-T611 and 7075-T612 in laboratory air and that of 7N01 conducted in dry air (dew point: − 70°C; water content: 2 ppm) at R = 0.513 are also plotted. In terms of ⌬Keff, no difference in crack growth rate between the T–L1 and T–L2 specimens was observed: the small difference in the ratio of foreign phase could be compensated, using effective stress intensity. Secondly, the crack growth rate in the present P/M Al alloy was the highest. This indicates that the crack growth rate of the present P/M alloy was inferior to other Al alloys. One reason is that the alloy was very brittle, and had a low fracture
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Figure 11 Crack growth rate, da/dN, as a function of ⌬Keff for T– L crack plane orientation
toughness: to improve the fatigue crack growth property, an increased ductility must be attained. Figure 12 illustrates SEM photographs of Type 3 fatigue fracture surfaces. At lower stress intensity, the crack grew with transgranular failure, that looked rather brittle (Figure 12(a)). With an increase in stress intensity, the amount of intergranular cracking increased (Figure 12(b)). At higher stress intensity, the crack grew with dimples (Figure 12(c)): an increase in fatigue crack growth rate may be explained by the presence of intergranular cracking as well as dimples due to low fracture toughness. Figures 13 and 14 illustrates the influence of crack plane orientation on fatigue crack growth rate of the present Al alloy, as a function of ⌬K and ⌬Keff, respectively. At lower stress intensity factor, there was little difference in crack growth rate between T–L and L–T crack plane orientation. However, with an increase in stress intensity factor, da/dN of T–L orientation was accelerated and was faster than that of L–T orientation, and the fatigue fracture toughness of L–T orientation was higher than T–L orientation. Figure 15 shows SEM photographs of the fatigue fracture surface of L–T orientation. Similar to a T–L specimen, fracture surface at lower stress intensity was dominated by transgranular cracking (Figure 15(a)); with an increase in stress intensity, brittle transgranular and intergranular cracking could be seen (Figure 15(b)), and the amount of dimples increased at higher stress intensity (Figure 15(c)). A most noticeable difference between T–L and L–T specimens is that a pull-out or fracture of intermetallic compound existed in L–T specimens. As was discussed earlier, intermetallic compound acted as a fiber reinforcement, and a reason why the fatigue crack growth rate in a L–T specimen was smaller than that of a T–L specimen
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Figure 13 Effect of crack plane orientation on fatigue crack growth rate in laboratory air as a function of ⌬K
Figure 12 Fatigue fracture surface of a T–L specimen conducted in laboratory air. (a) ⌬Keff ⬇ 1.5 MPa m1/2; (b) ⌬Keff ⬇ 2.8 MPa m1/2; (c) ⌬Keff ⬇ 3.7 MPa m1/2
can be partly explained by the prevention of crack growth by intermetallic compound aligned in the longitudinal direction.
Figure 14 Effect of crack plane orientation on fatigue crack growth rate in laboratory air as a function of ⌬Keff
Corrosion fatigue crack growth behavior Just after a fatigue pre-cracked specimen was exposed to a 3.5% NaCl solution, a crack grew for a while. However, with further stress cycles of about 104 (30 h at 0.1 Hz), the crack growth rate became
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was increased step by step, and finally the crack grew in a steady manner above ⌬Keff of 3 MPa m1/2. Figure 16 shows the crack growth rate under steady state as a function of ⌬Keff. The figure also illustrates the crack growth rates obtained before a decrease in growth rate and its consecutive crack retardation, which are plotted with a superscript *. The fatigue crack growth rate in a 3.5% NaCl solution was the same as that in laboratory air: the present P/M aluminum alloy was insensitive to NaCl solution as far as the present experiments are concerned. Under a steady state, the stress intensity was relatively large, and dimples could be seen and the fracture surface was similar to that conducted in air. This is consistent with no influence of 3.5% NaCl solution on fatigue crack growth being observed. Of course, further investigation into corrosion fatigue crack growth behavior at lower crack growth rate is required. CONCLUSIONS We evaluated the mechanical properties of a newly developed P/M aluminum alloy whose tensile strength reached about 800 MPa: tensile tests, slow strain rate tests (SSRTs) to evaluate the sensitivity against stress corrosion cracking, and fatigue strength as well as fatigue crack propagation in air and in a corrosive environment were investigated. The investigation yielded the following conclusions: 1. The developed P/M aluminum alloy has extremely high-tensile strength of about 800 MPa. However, elongation at the break was about 1%, and a speci-
Figure 15 Fatigue fracture surface of L–T specimen. (a) ⌬Keff ⬇ 1.7 MPa m1/2; (b) ⌬Keff ⬇ 2.7 MPa m1/2; (c) ⌬Keff ⬇ 3.5 MPa m1/2
smaller, and finally it was retarded. During this unsteady crack growth process, a crack opening stress intensity increased with stress cycles. The fracture surface where the crack was retarded was covered with thick corrosion products. These mean that an unsteady crack growth, or crack retardation was due to corrosion-product induced crack closure14,15. Because of this crack retardation, an applied load
Figure 16 Corrosion fatigue crack growth behavior of L–T specimens in a 3.5% NaCl solution. A superscript * indicates the crack growth rate was obtained before a decrease in growth rate and its consecutive crack retardation due to corrosion-products induced crack closure
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men fractured with little reduction in area and shear lip, showing its brittle nature. SSRT strength slightly decreases at lower displacement rate. However, the rate is smaller than that of other I/M aluminum alloys sensitive to stress corrosion cracking. The alloy has superior resistance to initiation of stress corrosion cracking. The fatigue crack growth rate of the alloy is higher than those of I/M Al–Zn–Mg alloys. One reason is that the alloy has a brittle nature, which is confirmed by a low fatigue fracture toughness of 7–8 MPa m1/2 and low elongation at the break. Another reason is that the fatigue fracture is associated with some amount of intergranular cracking. The fatigue crack growth rate in L–T orientation is lower than that in T–L orientation, because of an intermetallic compound (Al20(Cu,Zn)2Mn3) aligned in the longitudinal, or extruded, direction. Quasi-static tensile and fatigue crack is initiated at an inclusion: decreasing inclusions in the material is important for further improvement of its mechanical properties, including fatigue strength. The influence of 3.5% NaCl solution on fatigue crack growth is negligible as far as the present experiments are concerned. In NaCl solution, a crack is initiated at pitting corrosion on the surface. This promotes crack initiation, and therefore, the corrosion fatigue life of a smooth round specimen is lower than that conducted in air. In some cases a large pitting, or corrosion defect, occurs and SSRT strength as well as corrosion fatigue strength becomes extremely small. One of the most important future aims of developing a high-strength P/M aluminum alloy is to improve toughness, which in turn will improve fatigue strength as well as damage tolerance; the fabrication process must be optimized to avoid inclusions as
well as the foreign phases observed in some of the lots used. This will help the alloy to be improved in fatigue strength as well as corrosion properties. ACKNOWLEDGEMENTS The authors would like to thank Mr J. Kusui and Mr K. Yokoe, Research and Development Laboratory, Toyo Aluminium Company, Ltd., for the donation of the testing materials. REFERENCES 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15
Pickns, J. R., Jour. Mater. Sci., 1981, 16, 1437–1457. Mathy, A., Scharf, G., Becker, J., Fischer, G., Keinath, W., Gysler, A. and Lu¨tjering, G., Metall., 1990, 44, 532–539. Osamura, K., Kubota, O., Promstit, P., Okuda, H., Ochiai, S., Fujii, K., Kusui, J., Yokote, T. and Kubo, K., Metallurg. Trans., 1995, 26A, 1597–1599. Komai, K. and Minoshima, K., JSME International Journal, Series I, 1989, 32, 1–13. Komai, K. and Minoshima, K., Bulletin of Japan Society of Mechanical Engineers, 1985, 28, 2211–2218. Speidel, M. O., Metallurg. Trans., 1975, 6A, 631–651. Pickens, J. R. and Christodoulou, L., Metallurg. Trans., 1987, 18A, 135–149. Christodoulou, L., Gordon, J. R. and Pickens, J. R., Metallurg. Trans., 1985, 16A, 945–950. Metals Handbook, Vol. 2: Properties and Selection: Nonferrous Alloys and Pure Metals, ASM, Metals Park, Ohio, USA, 1979, p. 62. Ugiansky, G. M. and Payer, J. H., eds., Stress corrosion cracking—The slow strain-rate technique, ASTM STP 665, 1979. Minakawa, K., Levan, G. and McEvily, A. J., Metallurg. Trans., 1986, 17A, 1787–1795. Jono, M., Song, J. and Sugeta, A., J. Mater. Sci. Japan (in Japanese), 1985, 34, 1193–1199. Komai, K. and Minoshima, K., Trans. Japan Society of Mechanical Engineers (in Japanese), 1984, 50, 1804–1810. Endo, K., Komai, K. and Shikida, S., ASTM STP, 1984, 801, 81–95. Endo, K., Komai, K. and Ohnishi, K., Jour. Mater. Sci., Japan (in Japanese), 1968, 17, 160–168.
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Evaluation of stress corrosion resistance and corrosion fatigue fracture behavior of ultrahigh-strength P/M Al–Zn–Mg alloy Kohji Minoshima*‡, Makoto Okada† and Kenjiro Komai* *Department of Mechanical Engineering, Graduate School of Engineering, Kyoto University, Yoshida-Honmachi, Sakyo-Ku, Kyoto 606-8501, Japan †Central Japan Railway Co., Yaesu 1-6-6, Chuo-Ku, Tokyo 103-8288, Japan (Accepted 21 May 1998) Quasi-static tensile tests in air and slow strain rate tests (SSRTs) in a 3.5% NaCl solution were conducted in an ultra-high-strength P/M Al–Zn–Mg alloy fabricated through powder metallurgy. Attention is also paid to fatigue strength and fatigue crack growth behavior in laboratory air and in a 3.5% NaCl solution. The alloy has extremely high strength of about 800 MPa. However, elongation at break remains small, at about 1.3%. The final fracture occurs by a macroscopically flat crack normal to the tensile axis, with little reduction in area and little shear lip on the periphery of a smooth sample. However, it fails microscopically in a ductile manner, with dimples. Dimple size is less than 1 m, because the grain size of the alloy is extremely small. Strengthening mechanisms operating in the alloy are: small grains, sufficient metastable ⬘ phase in a matrix, and intermetallic compound acting as a fiber reinforcement. The SSRT strength in a 3.5% NaCl solution decreases slightly at a very low strain rate, that is smaller than those observed in aluminum alloys sensitive to stress corrosion. This means that the crack initiation resistance to stress corrosion is superior. However, under cyclic loading, the corrosion fatigue strength becomes lower than that conducted in air, because pitting corrosion on a sample surface acts as a stress concentrator. Crack initiation site of quasi-static and fatigue failure of the alloy is at inclusions, and hence, it is essential to decrease inclusions in the alloy for the improvement of the mechanical properties. Fatigue crack resistance of the alloy is inferior to conventional Al–Zn–Mg alloys fabricated by ingot metallurgy, because the fatigue fracture toughness, or ductility, of the alloy is inferior to other Al alloys, and intergranular cracking promotes crack growth. However, no influence of 3.5% NaCl solution on corrosion fatigue crack growth is observed, although an investigation is required into whether stress corrosion crack growth occurs or not, and at the same time, and of corrosion fatigue crack growth behavior at lower stress intensity. The fracture surface and crack initiation sites are closely examined using a high-resolution field emission type scanning electron microscope, and the fracture mechanisms of the alloy are discussed. 1998. Published by Elsevier Science Ltd. All rights reserved (Keywords: corrosion fatigue; stress corrosion cracking; SSRT; fatigue strength; fatigue crack propagation; fractography; powder metallurgy; Al–Zn–Mg alloy)
INTRODUCTION
and distribution of metastable precipitates acting as pinning centers for preventing dislocation motion. In particular, Mn and a very small amount of Ag additions greatly improve the tensile strength of the P/M alloys; the tensile strength exceeding 900 MPa3 is reported. However, little is known about its failure mechanism. Besides, in order to facilitate the application of these materials to machines and structures, the mechanical properties, including fatigue in various environments, have to be investigated4,5; the resistance of the materials to environments is extremely important. Important issues about mechanical properties in service are: quasi-static strength, ductility or fracture toughness, stress corrosion cracking (SCC) resistance, and also fatigue strength. The crack initiation resistance
Due to further progress in science and technology, new high-strength materials have become especially important; among these are metallic materials manufactured by powder metallurgy (P/M). One of the main attractions of P/M alloys is the ability to fabricate complex parts in an economical way: the ability of homogeneous microstructures and of element design of an alloy which is impossible by traditional ingot metallurgy (I/M). Recently, remarkable improvement in mechanical properties of P/M Al–Zn–Mg–Cu alloys has been achieved1–3 by controlling the morphology ‡Corresponding author.
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of stress corrosion cracking in a P/M Al alloy is reported to be superior to I/M aluminum alloy6,7, because a P/M alloy is subjected to rather general corrosion not to localized corrosion7. One reason is that large precipitates including Fe or Cu are not present in a P/M alloy, which act as cathodic sites to promote localized corrosion. However, the crack growth rate of a P/M alloy is almost the same as a traditional I/M Al alloy7,8. When these alloys are to be used in service, the properties in an operating environment under dynamic loading such as corrosion fatigue must be clarified. In this investigation, quasi-static tensile tests, slow strain rate tests (SSRTs) in a 3.5% NaCl solution, and fatigue strength and fatigue crack growth behavior in laboratory air and in a 3.5% NaCl solution were investigated in an ultra-high-strength P/M Al–Zn–Mg–Cu– Mn–Ag alloy3. Fracture surface and specimen surface after testing were closely examined using a field-emission type high-resolution scanning electron microscope, and the fracture mechanisms are discussed. EXPERIMENTAL PROCEDURES The material tested was a high-strength Al–9Zn–3Mg– 1.5Cu–4Mn–0.5Zr–0.04Ag alloy, fabricated by powder metallurgy3. The powders prepared by air atomizing were pressed into a rod-like shape by a cold isostatic pressing. The rod was first preheated in nitrogen gas at 500°C for 1 h, and extruded to an extrusion ratio of 20. The material was then machined to a specimen and heat-treated to a T6 condition: the specimens were solution-treated at 490°C for 2 h, then water-quenched, and aged at 120°C for 24 h. The shape and dimensions of test samples are shown in Figure 1: smooth round specimens shown in Figure 1(a) were machined in a LL orientation, where the loading axis paralleled the extrusion direction. To investigate crack growth behavior, compact tension
type specimens shown in Figure 1(b) were machined in T–L and L–T orientations. Four lots were used to prepare the specimens: see Table 1. Round smooth specimens (Lots 1, 2 and 3) were machined from a rod 40 mm in diameter, whereas CT specimens were machined from a square rod of 40 mm each side. A smooth sample surface was polished with #1500 emery paper followed by final finish with diamond paste of 2.5 m. Tensile tests were conducted in laboratory air at a displacement rate of 2 mm/min by using a computer-controlled tensile testing machine (Shimadzu Co. Ltd., Autograph AG-10TD, load capacity: 100 kN). To evaluate the SCC susceptibility, slow strain rate tests (SSRTs) were conducted at a displacement rate of 0.05 and 0.005 mm/min. A corrosive environment selected was a 3.5% NaCl solution prepared with reagent grade NaCl and deionized water (specific resistance > 1 M⍀ cm). The temperature of the solution was kept at 25°C by a thermoregulator, and was circulated by a vane pump made of synthetic resin between a corrosion chamber and a corrosion reservoir. In a 3.5% NaCl solution, a selected middle part of a sample of 2 mm in length was exposed to the environment, and the rest was coated to prevent the specimen surface from corrosion. Axial tension–tension fatigue tests at a stress ratio R of 0.1 were conducted on a smooth specimen in air and in a 3.5% NaCl solution at 25°C. In fatigue tests in air, a sinusoidal waveform at a stress cycle frequency f of 1–10 Hz was used, whereas in a 3.5% NaCl solution, a triangular wave at f = 0.1 Hz was selected. Similarly to in the SSRTs, a middle part of 2 mm in length was exposed to the environment. The testing machine employed was an electro-hydraulic fatigue testing machine with load capacity of 100 kN. Fatigue crack growth characteristics at a stress ratio R of 0.1 were evaluated by using a compact tension type specimen shown in Figure 1(b). In laboratory air, the effect of crack plane orientation on fatigue crack growth was investigated using T–L and L–T specimens. For L–T specimens, the influence of 3.5% NaCl solution was studied. In air, a sinusoidal waveform at f = 1–5 Hz was used, whereas a triangular waveform at f = 0.1 Hz was selected in a 3.5% NaCl solution. The testing machine employed was a computer-controlled electro-hydraulic fatigue testing machine (load capacity: 50 kN). Crack opening stress intensity was measured by an elastic unloading compliance technique: a strain gage was adhered to the back surface of a CT specimen, and a crack opening stress intensity was computed by using a personal computer linked with the fatigue testing machine. A surface crack length was monitored by using a traveling microscope. To compensate the tunneling effect, where a crack grew more in the inside
Table 1 P/M aluminum alloy for each set of experiments. Four lots were used to prepare the samples Smooth round specimens
Figure 1 Shape and dimensions of test specimens. All dimensions are in mm. (a) Smooth round specimen; (b) compact tension type specimen
CT specimens
Tensile tests in air
Lots 1, 2
SSRTs Fatigue tests in air Corrosion fatigue tests
Lot 1 Lots 1, 2, 3 Lots 1, 2 Lot 4
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of a specimen than on the surface, an average crack length was adopted to calculate a crack growth rate as well as stress intensity factor. An average crack length was defined by crack area divided by the sample thickness. EXPERIMENTAL RESULTS AND DISCUSSIONS Quasi-static tensile fracture behavior in air Table 2 summarizes the mechanical properties in laboratory air. The tensile strength went up to about 800 MPa which was slightly lower than the value reported in this material3. However, the strength was extremely high compared with ingot-metallurgy (I/M) 7XXX series aluminum alloys (B = 570–595 MPa: 7075-T6, 7175-T66)9. We must note that the elongation at the break was only 1.3%, which was smaller than I/M 7XXX series aluminum alloys (about 11% for 7075-T6 and 7175-T6)9. The fracture occurred by a macroscopic crack normal to the loading direction, and little shear lip could be seen (see Figure 2(a)): this is consistent with the alloy having small elongation at the break. The crack was initiated at an inside defect, a void or an inclusion (Figure 2(b)), whose diameter ranged from 70 to 100 m. This means that if the size of the defect could be made smaller, the mechanical properties might be increased. As shown in Figure 2(a), the alloy fractured in a brittle manner from the macroscopic view. However, the fracture surface was microscopically dominated by ductile fracture characterized by dimples (see Figure 2(c)). The strengthening mechanism of the alloy3 has been reported to be (1) small grain, less than 1 m, (2) metastable ⬘ phase in a matrix, and (3) intermetallic compound (Al20(Cu,Zn)2Mn3) that acts as a fiber reinforcement. The size of dimples observed on the fracture surface was about 0.5–1 m, and this suggested the strengthening mechanism due to small grain was achieved well. The other important strengthening mechanism that could be obtained from the fracture surface is intermetallic compound, which is shown by arrows in Figure 2(c). The intermetallic is about a few micrometers in length and the width was about a few hundred nanometers. The longitudinal direction of the intermetallic was parallel to the extruding direction, and they contributed to fiber reinforcement for the present P/M aluminum alloy. Slow-strain rate tests Figure 3 illustrates the fracture strength of SSRTs as a function of displacement rate in a 3.5% NaCl solution. The quasi-static fracture strengths in laboratory air are also plotted in the figure. Unlike the tensile test results in air, a crack that caused final failure in Table 2
Mechanical properties of the P/M aluminum alloy
Tensile strength (MPa) 797
Elongation (%) 1.31
Tensile tests were conducted in Lots 1 and 2, and two samples were used for each lot. The values listed in the table are the average values.
Figure 2 Fracture behavior of a quasi-static tensile test. (a) Macroscopic fracture morphology; (b) example of crack initiation site (inside inclusion); (c) fracture surface imaged with SEM. Arrows show examples of an intermetallic compound (Al20(Cu,Zn)2Mn3) that contributes to fiber reinforcement
a 3.5% NaCl solution was initiated at the sample surface. Figure 4 shows SEM photographs of a crack initiation site: a crack was initiated at a pitting corrosion site (Figure 4(a)), irrespective of loading condition, which will be refereed to as ‘pitting corrosion’. In some cases, severe corrosion occurred, and the size of the corrosion defect became still larger than pitting corrosion: this severe corrosion is presented in Figure 4(b). In Figure 3, symbols with a superscript * indicate the crack initiation site was such a large corrosion defect. When the crack initiation site was pitting corrosion, the strength in a 3.5% NaCl solution at 0.05 mm/min, which corresponded to 1.5–2.0 × 10−6 l/s,
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Figure 3 Fracture strength as a function of displacement rate in a 3.5% NaCl solution and in air. A superscript * indicates the crack initiation site was a large corrosion defect shown in Figure 4(b)
Figure 4 Crack initiation site of an SSRT in a 3.5% NaCl solution. (a) Pitting corrosion type (displacement rate: 0.005 mm/min, B = 732 MPa); (b) large corrosion defect type (displacement rate: 0.05 mm/min, B = 587 MPa)
was the same as that in laboratory air. Even in this case, the crack was initiated at a corrosion pit, not at an inside defect. If more tests were conducted at 0.05 mm/min, some inside defect could cause the crack initiation. However, at 0.005 mm/min corresponding to 1.5–2.0 × 10−7 l/s, the strength decreased from air data. In a material/environment system sensitive to stress
corrosion cracking, SSRT strength decreased with decreasing loading rate, and a minimum strength exists at a certain loading rate. Such a loading rate was reported to be 10−5 to 10−6 l/s10: this is larger than the loading rate of 0.005 mm/min (1.5–2 × 10−7 l/s) where the SSRT strength of the alloy decreased from air data. SSRTs evaluate the sensitivity of SCC in a given environment, in particular, the strength of crack initiation and passive films formed on a surface. When the crack initiation site was pitting corrosion, the strength slightly decreased at a lower loading rate of 0.005 mm/min. This decrease is small compared with that for materials sensitive to stress corrosion cracking: a decrease sometimes reaches almost half or smaller than half of the tensile strength10. This indicates that the alloy had a superior SCC strength to I/M aluminum alloys. When an aluminum alloy has very small grains such as this P/M alloy, large precipitates in a grain do not exist and, therefore, less damage of passive films is achieved, which in turn causes an increase in SCC strength. Fracture surface was dominated by dimples and corrosion products. However, at a loading rate of 0.005 mm/min, intergranular failure could be observed. The intergranular failure was typically observed in SCC of an aluminum alloy. Further investigation including crack growth of stress corrosion under a sustained load is required. When the crack initiation site was a large corrosion defect, the SSRT strength decreased from that conducted in air at both 0.05 and 0.005 mm/min, because a large corrosion defect increased the stress concentration, thereby decreasing the SSRT strength. What had caused such a severe corrosion, or corrosion defects, may be related to some kind of localized segregation due to insufficient control of the manufacturing process. Further investigation is necessary to avoid such severe corrosion: this will improve the properties of the alloy in a corrosive environment. Fatigue strength of smooth round specimens Figure 5 shows S–N curves under uniaxial tension– tension fatigue loading. The symbols with F indicate that a foreign phase could be seen in the fracture surface, which will be discussed later in detail. The fatigue results in air show a somewhat large scatter, and the specimens with foreign phases tended to have a lower fatigue strength than others. The macroscopic fracture morphology was almost the same as that fractured under quasi-static tensile loading in air: the specimen fractured with little shear lip on the periphery of the sample, and the fracture surface was normal to the axial loading direction. Under fatigue loading, a crack initiated at an inside inclusion near the specimen surface, irrespective of applied stress. Such a crack may be initiated by interfacial debonding between matrix and inclusion and/or fracture of the inclusion itself. The fracture surface consisted of fatigue and final catastrophic failure areas. The stress intensity factor obtained by using the crack length at final catastrophic failure was 7 to 8 MPa m1/2, that agreed with the fatigue fracture toughness of L– T orientation obtained by fatigue crack growth experiments, which will be shown later. Irrespective of applied stress, a fatigue crack was initiated at an inclusion near the specimen surface.
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Figure 5 S–N curves under uniaxial tension–tension fatigue loading. A superscript * indicates that the initiation site was an inclusion, while a superscript F indicates that a foreign phase could be seen in the fatigued fracture surface
Electron probe microanalysis showed that inclusions could be classified into three types: oxide inclusion, inclusion with foreign metals (Fe, Zn), and silica (SiO2). Inclusions with foreign metals or silica are considered to be brought about during manufacturing. Oxide inclusions are the most important and inherent problems for P/M metallurgy that we have to overcome. These results show that the amount of inclusions in the material should be decreased to improve fatigue strength of the present alloy. Fatigue fracture surface was similar to those observed in CT specimens shown in Figure 15. Near the crack initiation site, the fracture surface was transgranular, and macroscopically it looked flat, because the grain size of the material was extremely small. With an extension of fatigue crack, not only transgranular cracking but also intergranular cracking appeared. Fracture of intermetallic compound (Al20(Cu,Zn)2Mn3) could also be seen. At the final catastrophic failure stage, dimples as well as fracture of intermetallic compounds dominated over the fracture surface. Corrosion fatigue strength The corrosion fatigue results in a 3.5% NaCl solution are shown in Figure 5. A crack that caused final failure was initiated at either an inclusion or a corrosion pit shown in Figure 6(a). A superscript * is added to the symbol when the initiation site was an inclusion. When the corrosion fatigue life was extremely short at 300 MPa (Nf = 5.83 × 102), an extremely large corrosion-induced defect of 500 to 600 m was an initiation site: see Figure 6(b). This large corrosion defect caused a large decrease in corrosion fatigue life compared to other specimens. Except for this, the fatigue strength in a 3.5% NaCl solution was even smaller than that in air. As is discussed later, little influence of the corrosion environment on fatigue crack growth in the alloy was observed. Therefore, a decrease in corrosion fatigue
Figure 6 Crack initiation site of a corrosion fatigue test in a 3.5% NaCl solution. (R = 0.1, f = 0.1 Hz). (a) Pitting corrosion (max = 300 MPa, Nf = 1.51 × 104); (b) large corrosion defect (max = 300 MPa, Nf = 5.83 × 102)
life was due to promotion of crack initiation by pitting corrosion. Corrosion pits tended to occur along the extruding direction, and these pits coalesced with each other, resulting in a crack initiation site. In some cases, a large corrosion defect was formed on a specimen surface, shown in Figure 4(b) and Figure 6(b), resulting in a further decrease in SSRT strength or corrosion fatigue life. A reason why such large corrosion defects occurred may be due to some localized segregation of constituents: further investigation is required to improve the corrosion property which in turn improves the mechanical properties of the P/M alloy in a corrosive environment. Fatigue crack growth in air Figure 7 shows a macroscopic fatigue fracture surface of a CT specimen conducted in air: three different types of fracture morphology could be seen. These will be refered to as Type 1, Type 2, and Type 3, respectively. Respective microscopic fracture surfaces are shown in Figure 8, and the characteristic features of each fracture surface are summarized as follows: Type 1 fracture surface looked dark, and the fracture surface was completely different from the other types. In the case of L–T specimens, a crack tended to be retarded when the crack reached the Type 1 fracture surface. Type 2 fracture surface involves some typical fracture surface morphology which looked as if the mating surfaces had contacted each other. As for Type 3 fracture surface shown in Figure 8(c) at higher stress range, characteristic fracture morphology of this P/M alloy, small grains and intermetallic compound (Al20(Cu,Zn)2Mn3) acting as a reinforcement, could be seen. Type 1 and Type 2 fracture surfaces were in
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Figure 7 Macroscopic fatigued fracture surface of T–L orientation (CT specimen)
particular observed in Lot 4, that was machined into CT specimens. The different fracture morphology was due to the existence of foreign phases in the material which were elongated in the extrusion direction: the foreign phases are considered to be caused by insufficient control of a manufacturing parameter during extrusion. A ratio of each fracture surface morphology, or foreign phases, depended on a specimen and, therefore, crack growth rate varied from sample to sample. Hence, in this experiment, crack growth rates obtained by a CT specimen, of which a Type 3 fracture surface dominated over the fracture surface, are plotted. Figure 9 illustrates the fatigue crack growth rate of T–L crack plane orientation in air as a function of ⌬K. The crack growth rates of an I/M Al–Zn–Mg alloy, 7075-T611,12 and P/M Al–Zn–Mg alloy, 7090-T611 are also plotted for comparison. At lower stress intensity, there was little difference in fatigue crack growth rate. At an intermediate ⌬K, crack growth rate of T–L2 was accelerated. However, at higher stress intensity, acceleration of crack growth was observed in both T–L1 and T– L2 specimens. The fatigue fracture toughness of the alloy was about 6–7 MPa m1/2, and the crack growth rate increased as a maximum stress intensity reached the fatigue fracture toughness. We must note that crack growth rate at higher stress intensity of the present alloy was the highest, and decreased in the order of 7090-T6 and 7075-T6. This means that the fatigue fracture toughness of P/M Al alloys, in particular the present Al alloy, was inferior to conventional I/M Al alloys. Figure 10 shows the crack opening stress intensity factor as a function of maximum stress intensity factor.
Figure 8 Microscopic fatigued fracture surface of a CT specimen conducted in air (T–L orientation); (a) Type 1; (b) Type 2; (c) Type 3
The crack opening stress intensity was higher in the T–L1 specimen than the T–L2 specimen. The ratio of foreign phases of T–L1 was higher than that of the T–L2 specimen, and the roughness of fracture surface was higher in T–L1 specimen. This indicates that the crack opening stress intensity increased by roughness induced crack closure. If the ratio of foreign phases of T–L1 specimens was smaller, the crack opening stress intensity factor would be decreased.
Stress corrosion resistance and corrosion fatigue fracture behavior of P/M Al–Zn–Mg alloy
Figure 9 Crack growth rate, da/dN, as a function of ⌬K for T–L crack plane orientation
Figure 10 Crack opening stress intensity factor as a function of maximum stress intensity factor
The crack growth rates as a function of effective stress intensity factor, ⌬Keff, are shown in Figure 11. In the figure, crack growth rates of 7090-T611 and 7075-T612 in laboratory air and that of 7N01 conducted in dry air (dew point: − 70°C; water content: 2 ppm) at R = 0.513 are also plotted. In terms of ⌬Keff, no difference in crack growth rate between the T–L1 and T–L2 specimens was observed: the small difference in the ratio of foreign phase could be compensated, using effective stress intensity. Secondly, the crack growth rate in the present P/M Al alloy was the highest. This indicates that the crack growth rate of the present P/M alloy was inferior to other Al alloys. One reason is that the alloy was very brittle, and had a low fracture
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Figure 11 Crack growth rate, da/dN, as a function of ⌬Keff for T– L crack plane orientation
toughness: to improve the fatigue crack growth property, an increased ductility must be attained. Figure 12 illustrates SEM photographs of Type 3 fatigue fracture surfaces. At lower stress intensity, the crack grew with transgranular failure, that looked rather brittle (Figure 12(a)). With an increase in stress intensity, the amount of intergranular cracking increased (Figure 12(b)). At higher stress intensity, the crack grew with dimples (Figure 12(c)): an increase in fatigue crack growth rate may be explained by the presence of intergranular cracking as well as dimples due to low fracture toughness. Figures 13 and 14 illustrates the influence of crack plane orientation on fatigue crack growth rate of the present Al alloy, as a function of ⌬K and ⌬Keff, respectively. At lower stress intensity factor, there was little difference in crack growth rate between T–L and L–T crack plane orientation. However, with an increase in stress intensity factor, da/dN of T–L orientation was accelerated and was faster than that of L–T orientation, and the fatigue fracture toughness of L–T orientation was higher than T–L orientation. Figure 15 shows SEM photographs of the fatigue fracture surface of L–T orientation. Similar to a T–L specimen, fracture surface at lower stress intensity was dominated by transgranular cracking (Figure 15(a)); with an increase in stress intensity, brittle transgranular and intergranular cracking could be seen (Figure 15(b)), and the amount of dimples increased at higher stress intensity (Figure 15(c)). A most noticeable difference between T–L and L–T specimens is that a pull-out or fracture of intermetallic compound existed in L–T specimens. As was discussed earlier, intermetallic compound acted as a fiber reinforcement, and a reason why the fatigue crack growth rate in a L–T specimen was smaller than that of a T–L specimen
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Figure 13 Effect of crack plane orientation on fatigue crack growth rate in laboratory air as a function of ⌬K
Figure 12 Fatigue fracture surface of a T–L specimen conducted in laboratory air. (a) ⌬Keff ⬇ 1.5 MPa m1/2; (b) ⌬Keff ⬇ 2.8 MPa m1/2; (c) ⌬Keff ⬇ 3.7 MPa m1/2
can be partly explained by the prevention of crack growth by intermetallic compound aligned in the longitudinal direction.
Figure 14 Effect of crack plane orientation on fatigue crack growth rate in laboratory air as a function of ⌬Keff
Corrosion fatigue crack growth behavior Just after a fatigue pre-cracked specimen was exposed to a 3.5% NaCl solution, a crack grew for a while. However, with further stress cycles of about 104 (30 h at 0.1 Hz), the crack growth rate became
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was increased step by step, and finally the crack grew in a steady manner above ⌬Keff of 3 MPa m1/2. Figure 16 shows the crack growth rate under steady state as a function of ⌬Keff. The figure also illustrates the crack growth rates obtained before a decrease in growth rate and its consecutive crack retardation, which are plotted with a superscript *. The fatigue crack growth rate in a 3.5% NaCl solution was the same as that in laboratory air: the present P/M aluminum alloy was insensitive to NaCl solution as far as the present experiments are concerned. Under a steady state, the stress intensity was relatively large, and dimples could be seen and the fracture surface was similar to that conducted in air. This is consistent with no influence of 3.5% NaCl solution on fatigue crack growth being observed. Of course, further investigation into corrosion fatigue crack growth behavior at lower crack growth rate is required. CONCLUSIONS We evaluated the mechanical properties of a newly developed P/M aluminum alloy whose tensile strength reached about 800 MPa: tensile tests, slow strain rate tests (SSRTs) to evaluate the sensitivity against stress corrosion cracking, and fatigue strength as well as fatigue crack propagation in air and in a corrosive environment were investigated. The investigation yielded the following conclusions: 1. The developed P/M aluminum alloy has extremely high-tensile strength of about 800 MPa. However, elongation at the break was about 1%, and a speci-
Figure 15 Fatigue fracture surface of L–T specimen. (a) ⌬Keff ⬇ 1.7 MPa m1/2; (b) ⌬Keff ⬇ 2.7 MPa m1/2; (c) ⌬Keff ⬇ 3.5 MPa m1/2
smaller, and finally it was retarded. During this unsteady crack growth process, a crack opening stress intensity increased with stress cycles. The fracture surface where the crack was retarded was covered with thick corrosion products. These mean that an unsteady crack growth, or crack retardation was due to corrosion-product induced crack closure14,15. Because of this crack retardation, an applied load
Figure 16 Corrosion fatigue crack growth behavior of L–T specimens in a 3.5% NaCl solution. A superscript * indicates the crack growth rate was obtained before a decrease in growth rate and its consecutive crack retardation due to corrosion-products induced crack closure
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men fractured with little reduction in area and shear lip, showing its brittle nature. SSRT strength slightly decreases at lower displacement rate. However, the rate is smaller than that of other I/M aluminum alloys sensitive to stress corrosion cracking. The alloy has superior resistance to initiation of stress corrosion cracking. The fatigue crack growth rate of the alloy is higher than those of I/M Al–Zn–Mg alloys. One reason is that the alloy has a brittle nature, which is confirmed by a low fatigue fracture toughness of 7–8 MPa m1/2 and low elongation at the break. Another reason is that the fatigue fracture is associated with some amount of intergranular cracking. The fatigue crack growth rate in L–T orientation is lower than that in T–L orientation, because of an intermetallic compound (Al20(Cu,Zn)2Mn3) aligned in the longitudinal, or extruded, direction. Quasi-static tensile and fatigue crack is initiated at an inclusion: decreasing inclusions in the material is important for further improvement of its mechanical properties, including fatigue strength. The influence of 3.5% NaCl solution on fatigue crack growth is negligible as far as the present experiments are concerned. In NaCl solution, a crack is initiated at pitting corrosion on the surface. This promotes crack initiation, and therefore, the corrosion fatigue life of a smooth round specimen is lower than that conducted in air. In some cases a large pitting, or corrosion defect, occurs and SSRT strength as well as corrosion fatigue strength becomes extremely small. One of the most important future aims of developing a high-strength P/M aluminum alloy is to improve toughness, which in turn will improve fatigue strength as well as damage tolerance; the fabrication process must be optimized to avoid inclusions as
well as the foreign phases observed in some of the lots used. This will help the alloy to be improved in fatigue strength as well as corrosion properties. ACKNOWLEDGEMENTS The authors would like to thank Mr J. Kusui and Mr K. Yokoe, Research and Development Laboratory, Toyo Aluminium Company, Ltd., for the donation of the testing materials. REFERENCES 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15
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