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Enhancement of the fatigue properties in a weldable high-strength Al-Zn-Mg-Mn alloy by means of Mn-dispersoids
International Journal of Fatigue 21 (1999) 383–391
Enhancement of the fatigue properties in a weldable high-strength Al-Zn-Mg-Mn alloy by means of Mn-dispersoids Duck Hee Lee a, Kyu Cheol Kim b, Dong Seok Park c, Soo Woo Nam
d,*
a
d
Department of Mat. Sci. and Eng., Korea Advanced Institute of Science and Technology, 373-1 Kusong-dong Yusong-gu, Taejon 305-701, South Korea b Andersen Consulting, 17-3 Yoido-dong Youngdeoungpo-ku, Seoul 150-010, South Korea c Defense Quality Assurance Agency, 341 Samjeongja-dong, Changwon 641-160, South Korea Dept. of Mat. Sci. and Eng., Korea Advanced Institute of Science and Technology, 373-1 Kusong-dong Yusong-gu, Taejon 305-701, South Korea Received 10 August 1998; received in revised form 29 October 1998; accepted 19 November 1998
Abstract This present research was undertaken to investigate the effects of Mn dispersoids on low cycle fatigue (LCF) and high cycle fatigue (HCF) in the peak aged Al-Zn-Mg-Mn alloy. The Mn addition in an Al-Zn-Mg-Mn alloy was found to form Mn dispersoids in sizes ranging from 0.05 to 0.5 µm. Mn dispersoids in an Al-Zn-Mg-Mn alloy results in increasing strength without sacrificing ductility by its dispersion hardening effects but at the same time homogeneous deformation. The homogeneous deformation is caused by the cross-slip of dislocations at Mn dispersoids and prevents stress and/or strain concentrations. In this Mn dispersoid containing alloy, both LCF and HCF lives are found to be significantly extended compared with that without any Mn dispersoids. This phenomenon is very interesting and is considered to be due to the simultaneous improvement of strength and ductility by Mn dispersoids. In this investigation, the reason why these excellent properties are observed is discussed in terms of the interaction between Mn dispersoids and the dislocation. 1999 Elsevier Science Ltd. All rights reserved. Keywords: Al-Zn-Mg-Mn alloy; Low cycle fatigue; Fatigue crack growth rate; Mn dispersoids; Cross-slip
1. Introduction The typical commercial 7XXX aluminum alloys are divided into two groups by the characteristic properties, i.e., the one is the weldable aluminum alloys of 7039 and 7017 etc. (UTS: 440–480MPa) which are being produced as thick plate to be used for the structural materials of transportation, aeroplanes, and military vehicles. The other is the non-weldable high strength aluminum alloys of 7075 and 7050 etc. (UTS: 540– 570MPa) which are being produced by extrusion process and these alloys have about 20–25% higher strength compared to that of the weldable alloys. It is well known that a small amount of transient elements such as Zr, Cr, and Mn form various dispersoids which affect the mechanical properties of aluminum
* Corresponding author. Tel.: +82 42 869 3318; fax: +82 42 869 3310; e-mail: [email protected].
alloys.[1–5] With this basic concept, adding more than 0.5wt.% Mn to the weldable Al-Zn-Mg series aluminum alloy, a new weldable high-strength aluminum alloy is developed.[6,7] This new weldable high-strength aluminum alloy has strength as high as that of Al 7075 and at the same time it has good weldability, as good as Al 7039. Therefore, this new alloy can be used to produce many structural parts of transportation, high speed trains, aeroplanes, and military vehicles. Mn dispersoids in a new Al-Zn-Mg-Mn alloy are found to act as the major agents in increasing strength without sacrificing ductility.[8,9] They also do not affect the weldability of this alloy, as well as the kinetics of η⬘ precipitation during the artificial aging process.[8] In general, the fatigue is classified by two regimes, low-cycle fatigue (LCF) and high-cycle fatigue (HCF). These two different fatigue phenomena have been known to be governed by different mechanical properties, i.e., LCF is mainly related with ductility and HCF is strongly affected by strength. Unfortunately, it is gen-
0142-1123/99/$ - see front matter 1999 Elsevier Science Ltd. All rights reserved. PII: S 0 1 4 2 - 1 1 2 3 ( 9 8 ) 0 0 0 8 5 - 1
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D. H. Lee et al. / International Journal of Fatigue 21 (1999) 383–391
erally accepted that, in metals, increase of strength results in the reduction of ductility. Therefore, it has been reported that commercial aluminum alloys with good LCF life have poor HCF strength and vice versa. In the present study, the role of Mn dispersoids in LCF and HCF behavior of Al-Zn-Mg-Mn alloys is investigated through TEM analysis in combination with fatigue tests.
tensile test machine attached to a tension-compression auto-toggling controlling system. These SEM in-situ fatigue tests were conducted at a strain rate of 1 mm/min and a total strain range of ±2%. Fatigue test specimens were solution treated at 460°C for 2 h, water quenched, and then tested.
3. Results and discussion 2. Experimental procedures
3.1. Microstructure and tensile properties
The chemical compositions of the alloys tested are presented in Table 1. The Al-Zn-Mg-(Mn) alloys were prepared by melting 99.99% pure Al with the appropriate master alloys and casting in an inert atmosphere. After homogenization treatment at 460°C for 24 h which is to form the Mn dispersoids and homogenize the segregation, the billets were extruded at 400°C to be 25 mm × 100 mm plates. Hot rolling of billets to 25 mm × 150 mm plates was also performed at 390°C with a final reduction ratio of 80%. The plates were solution treated at 460°C for 90 min, followed by a water quench. A twostep artificial aging treatment was carried out at 97°C for 8 h and then at 140°C for 10 h to give higher strength and less effective quench sensitivity, compared to the conventional one-step aging treatment. Low-cycle fatigue test specimens which had a gauge length of 6 mm and a gauge diameter of 4.5 mm were machined from the extruded and rolled plates with the stress axis parallel to the extrusion and rolling direction. The fully reversed axial strain-controlled low-cycle fatigue tests were conducted in air using an Instron 1380 with a constant strain rate of 4×10⫺3s⫺1. Fatigue crack growth tests were performed using a 10 ton capacity Instron 8500 dynamic machine at room temperature in air environment, with the procedures outlined in ASTM E647-88. Specimens were cut from the extruded and rolled plates in the LT orientation and machined to be the standard compact tension specimens with chevron notch. All the tests were carried out with 5 Hz triangular frequency under the load ratio of 0.25. Crack length was measured during test using a traveling microscope. For the SEM in-situ observation of the deforming surface morphology, in-situ low-cycle fatigue test specimens were made with a gauge length of 5.5 mm and a diameter of 2.5 mm, and cycled in a JEOL SM TS40
The morphology and distribution of Mn-dispersoids within two different contents of manganese are shown in Fig. 1. In the Mn bearing alloy, with an extruded and rolled Mn08 alloy, spherical- and rod-shaped Mn-dispersoids of the size in range from 0.05–0.5 µm are found to be dispersed uniformly in the matrix. The distribution and morphology of Mn dispersoids in rolled Mn00 and Mn08 alloys are almost consistent with those of extruded ones, respectively, even though the results of rolled alloys are not presented in this paper. It has been experimentally observed that the Mn-dispersoids were formed during a homogenization treatment and no change is
Table 1 Chemical composition of the investigated Al-Zn-Mg-(Mn) alloys(in wt.%) Alloy
Zn
Mg
Mn
Zr
Al
Mn00 Mn08
4.2 4.3
2.6 2.8
0.07 0.8
0.15 0.15
bal. bal.
Fig. 1. TEM micrographs showing the size and distribution of Mn dispersoids in the extruded alloys: (a) Mn00 alloy, (b) Mn08 alloy.
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found in the morphology and size during the aging treatment.[8] Also, according to the observation of Park,[8] Mn-dispersoids in the Al-Zn-Mg-Mn alloy is found not to affect the kinetics of precipitation by the artificial aging process. Mechanical properties of the investigated alloys are tabulated in Table 2. Both the yield and ultimate tensile strength of these Mn bearing alloys are measured to be increased with sufficient amount of elongation. This result implies that the fine Mn dispersoids in the matrix could act as a restriction for any dislocation motion, and consequently increase the strength of the Mn bearing alloy, extruded and rolled. 3.2. Tensile deformation behavior In order to identify the deformation mode of the investigated Al alloys, a dislocation structure was investigated using the specimen strained up to 1% as shown in Fig. 2. The TEM micrographs in Fig. 2 show that there is a trend of slip homogenization with increasing the density of Mn-dispersoids. It appears that this alternation in deformation behavior is due to the Mn dispersoids, since no significant difference in microstructure was found except this distribution of the Mn dispersoids. In Fig. 2(a), Mn00 alloy shows the wide width of intense slip bands consisting of high density of dislocations. However, Mn-dispersoids disperse slip and inhibit the formation of intense slip bands as shown in Fig. 2(b). Therefore, plastic deformation is more homogeneous. In both extruded and rolled conditions, it is believed that the dispersive distribution of Mn-dispersoids in Mn08 alloy promotes homogeneous deformation by converting the deformation mechanism from planar slip, which is well developed in the Mn00 alloy with no Mn dispersoids, to wavy slip. As for the homogeneous deformation in Mn08 alloy, the deformation mode could be mainly dominated by cross slip.[10] To identify the interaction behavior between Mn dispersoid and dislocation, TEM analysis was introduced with the specimen imposed small amount of plastic deformation which will be in detail discussed in the ‘TEM analysis’ section (Fig. 7). From the above TEM observation and previous works,[6,9] the role of Mn dispersoids in the Mn08 alloy is very clearly seen in that they act as a strengthener to block the motion of Table 2 Tensile properties Alloy
Condition
YS (MPa)
UTS (MPa) El. (%)
Mn00
Extruded Rolled Extruded Rolled
420 375 484 408
463 445 545 468
Mn08
15.2 13.0 12.3 14.3
Fig. 2. Comparison of the distribution of slip band after 1% tensile strain in the extruded alloys: (a) Mn00 alloy, B=[111], g=⬍220⬎; (b) Mn08 alloy, B=[111], g=⬍220⬎.
dislocations, and at the same time they distribute the slip of dislocations uniformly so as to enhance ductility. 3.3. Low cycle fatigue A Coffin-Manson plot is shown in Fig. 3. In this figure, we can easily see that the Mn bearing alloy, extruded and rolled, has a longer LCF life than no Mn bearing alloy in both extruded and rolled conditions, even though the strength of the Mn08 alloy is higher than that of the Mn00 alloy. Usually, a metal having higher elongation or ductility with lower strength has a longer LCF life than a low-elongation metal with high strength because the LCF test mode is strain-controlled. However, the Mn08 alloy shows sufficient ductility with higher strength but also has longer LCF life, compared with the Mn00 alloy. Cyclic properties of the investigated alloys are tabulated in Table 3. In order to explain the above mentioned results of the LCF life, the dislocation structure was investigated using the specimens fatigued up to 10% of the LCF life, as shown in Fig. 4. The result in Fig. 4 is similar to that observed in a previous section which was obtained
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Fig. 3.
Coffin-Manson plot of each alloy.
Table 3 Cyclic properties of low cycle fatigue specimens (Nfα·⌬⑀p = C, C = constant) Alloy
Condition
C
⫺a
Mn00
Extruded Rolled Extruded Rolled
0.188 0.239 0.263 0.402
0.667 0.670 0.688 0.728
Mn08
Fig. 4. TEM micrographs showing the dislocation arrangement formed after cyclic deformation in the rolled alloys: (a) Mn00 alloy, (b) Mn08 alloy.
throughout tensile deformation. When the inhomogeneous deformation occurs by planar slip in the Mn00 alloy, as shown in Fig. 4(a), the stress concentration around the pile-up dislocations at the end of slip band may occur and interact with grain boundaries. This local stress concentration may cause early fatigue crack initiation at grain boundaries which would then decrease LCF life. On the other hand, since the Mn dispersoids make the deformation homogeneous, as shown in Fig. 4(b), fatigue crack initiation due to the local stress concentration at grain boundaries is suppressed in the Mn bearing Mn08 alloy. Therefore, it appears that homogeneous deformation by means of the Mn dispersoids will lead to a longer LCF life. These phenomena will be in detail discussed in the following section. 3.4. SEM in-situ fatigue test The results of in-situ fatigue tests of the Mn00 alloy show typically coarse slip bands formed at 45° angles to the stress axis and often concentrated at grain boundary areas, as shown in Fig. 5. The slip band traces impinged at the grain boundaries may build up local stresses that can cause fatigue crack initiation. These
cracks are continuously extended and coalesce together during further cycling. In Fig. 5(c), the enlarged grain boundary cracks after 1100 cycles are shown. The localized slip bands in this alloy without Mn were found to be caused by a planar dislocation movement (Fig. 4(a)). The corresponding results of in-situ fatigue testing of the Mn08 alloy are shown in Fig. 6. This alloy containing Mn-dispersoids does not show any well-developed slip bands throughout its fatigue life. Thus it is believed that this kind of deformation mode is caused by the Mndispersoids. They omit coarse slip band formation and cause homogeneous deformation. As the number of fatigue cycles increases, small cracks are seen all over the specimen’s surface. These cracks are distributed, not at grain boundaries, but inside the grains. This indicates that they were neither caused nor affected by the concentrated slip bands close to the grain boundaries. This homogeneous deformation of the Mn08 alloy makes crack initiation difficult and extends its fatigue life.
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Fig. 5. SEM micrographs from in-situ fatigue tests of the Mn00 alloy at ⌬⑀t=±2%. Arrows indicate fatigue stress axis: (a) after 120 cycles, (b) after 790 cycles (c) after 1100 cycles.
3.5. TEM analysis To verify the cause of homogeneous deformation mode by Mn-dispersoids during fatigue testing, TEM analysis on the dislocation structure was conducted for the Mn08 alloy. The observations were made after only one cycle to avoid a very complicated dislocation structure due to multiplication and tangling. The TEM micrographs and corresponding schematic diagrams, which show the 3-D configuration of dislocations, are shown in Fig. 7. In order to analyze the dislocation, before and after its interaction with the Mn-dispersoids during fatigue testing, trace analysis was applied.[11,12] Obtaining the Burgers vector, the usual g·b criterion was used, and for the line direction determination, 3 beam directions were chosen to get the diffraction pattern. Using the relationship between the projected dislocation and that of the real, the unknown value of the real dislocation direction can be obtained from the known projected one.[11] The slip plane of the dislocations, containing both Burgers vector and dislocation line vector, can be easily obtained by the cross product of the two
vectors. Calculated Burgers vector, line vector and slip plane are listed in Table 4. In Fig. 7, dislocations are found to be moving on a (111) slip plane (a, b and c) until they are blocked by the non-shearing Mn-dispersoids and the dislocations tend to cross-slip to (111) slip plane (d and e). This indicates that dislocations crossslip from primary slip plane, (111), to the other active slip plane, (111), by the Mn-dispersoids without any dislocation pile-up at the Mn-dispersoid. Such a phenomenon of blocking and cross-slip of dislocations corresponds with the SEM in-situ test results in that the slip steps are very fine and evenly distributed. This suggests slip homogenization through the cross-slip over the nonshearable obstacles. Therefore, it is believed that the dispersive distribution of the coarser Mn-dispersoids in a Mn08 alloy promotes homogeneous deformation by converting the dislocation movement from planar slip to a wavy slip mode through cross-slip. This results in increasing strength without sacrificing ductility, as well as retarding fatigue crack initiation and propagation. All of this is generated by the inhomogeneous deformation within the grain boundary.
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Fig. 6. SEM micrographs from in-situ fatigue tests of the Mn08 alloy at ⌬⑀t=±2%. Arrows indicate fatigue stress axis: (a) after 170 cycles, (b) after 800 cycles (c) after 1200 cycles.
Fig. 7. TEM micrograph showing the dislocation structure of the Mn08 alloy after 1 cycle at ⌬⑀t=±1.7%. (B=[112], g=[220]; Arrows indicate Mn-dispersoids).
D. H. Lee et al. / International Journal of Fatigue 21 (1999) 383–391
Table 4 Dislocation characters and slip planes in the Mn08 alloy Dislocation b a b c d e
1/2[0 1/2[0 1/2[0 1/2[0 1/2[0
1¯ 1¯ 1¯ 1¯ 1¯
1] 1] 1] 1] 1]
x
Slip plane
Slip system
[61 50 11] [152 117 35] [9 15 6¯] [23 8 15] [9 16 7]
(1¯ (1¯ (1¯ (1 (1
primary primary primary secondary secondary
1 1 1 1 1
1) 1) 1) 1) 1)
3.6. High-cycle fatigue For the high-cycle fatigue test, the fatigue crack growth rate has been measured and plotted in terms of ⌬K, as shown in Fig. 8. In Fig. 8, the threshold stress intensity (⌬Kth) values were obtained from the plot of da/dN vs. ⌬K within the region I suggested by Ritchie et al.[13] The value of ⌬Kth is considered to strongly depend on alloy composition, ingot process, thermomechanical treatment, test condition and so on.[14–17] The fatigue crack growth tests at several different conditions including that presented in this article were conducted using the Mn00 and Mn08 alloys. Under all test conditions, it was found that threshold values occurred at about 10⫺7 m/cycle.[14,15] Therefore, even though the ⌬Kth values in the present study are thought to be somewhat higher than those of commercial Al alloys, this finding seems to be characteristic of these tested alloys. In addition, from the literature survey,[17, 18 (pp. 519– 618), 19 (pp. 56–104)] it can be seen that as the threshold value of a material is higher, the threshold value
389
generally occurs at higher growth rate (higher da/dN) although this event is not always true. The Mn08 alloy is found to have a larger threshold stress intensity than the Mn00 alloy in both extruded and rolled conditions. Also, at the same stress intensity which was over the threshold stress intensity, the Mn08 alloy has a much slower crack growth rate compared to that of the Mn00 alloy and thus, the fatigue crack growth resistance of the Mn08 alloys is superior to those of the Mn00 alloys. These results are thought to be due to Mn dispersoids, since the major difference between Mn bearing and none Mn containing alloys is whether there are Mn dispersoids or not. The fatigue crack growth properties of these investigated aluminum alloys are summarized in Table 5. On the other hand, the difference in fatigue crack growth properties between the extruded and rolled alloys is thought to be due to the aging treatments which are considered to be a major affecting factor and grain structures, because these are the main differences for two kinds of alloys with the different treatments (extrusion and rolling). The different aging treatment was conducted on extruded specimens for the fatigue crack growth test (Fig. 8). Extruded alloys (Mn00 and Mn08 alloy) have the partially recrystallized grains within the elongated fibrous grains along the extrusion direction. As the content of Mn increases, the recrystallized grain size somewhat decreases. The grain sizes of Mn00 and Mn08 alloys are about 13 µm and 8 µm, respectively. However, the rolled alloys have the coarser pancake shaped grain structure (⬎30 µm) than those of the extruded alloys with same Mn composition. To identify the above mentioned results, the fracture surfaces in extruded alloys were examined by SEM. Fig. 9 shows the differences in the morphology of the fracture surfaces. The extruded Mn08 alloy has a much more rough fracture surface. Therefore, in the Mn08 alloy, it is supposed that as the fatigue crack advances, the Mn dispersoids may cause crack deflections and make the crack path more tortuous by acting as an obstacle to the advance of a fatigue crack. These results illustrate that Mn dispersoids can increase many more deflections in the fatigue crack path than the inhomogeneous deformation, causing the fatigue crack propagation curve to shift from the left to right, as illustrated in Fig. 8. Table 5 Fatigue crack propagation properties of the investigated Al alloys (da/dN = A(⌬K)m)
Fig. 8. Influence of the Mn content on the fatigue crack propagation behaviors. (Extruded alloys: natural aged at R.T./96h → 100°C/10 min +160°C/180 min; YS(MPa)/UTS(MPa)/El.(%) =258/344/20 (Mn00 alloy), 360/428/16.5 (Mn08 alloy)).
Alloy
Condition
⌬Kth (MPa·m1/2)
A m (×10⫺11)
Mn00
Extruded Rolled Extruded Rolled
8.22 15.46 10.18 18.38
1.70 1.75 1.92 0.06
Mn08
3.74 3.59 3.68 4.32
390
D. H. Lee et al. / International Journal of Fatigue 21 (1999) 383–391
The Mn dispersoids in the Mn08 alloy increase the tortuosity of the crack path and improve the fatigue crack growth properties, through a crack deflection caused by the finely dispersed Mn dispersoids. Considering the results of the LCF and HCF tests in the extruded and rolled Mn bearing alloy, it can be seen that finely dispersed Mn dispersoids can improve not only the LCF life but also resistance to fatigue crack growth.
Acknowledgements This work was supported by MOST in the regime of the HAN (NM05630) project. Thanks are given for its financial support.
References
Fig. 9. SEM micrographs showing the fatigue crack growth path and fracture surface in the extruded alloys: (a) Mn00 alloy (b) Mn08 alloy.
The structural components have notches of different acuity. In the view point of notch sensitivity of a material, the sharp notch may cause earlier fatigue short crack initiation at the notch tip. However, once the fatigue short crack is, to some extent, extended and not affected by stress field related with notch, the sharpness of notch may have less effect on crack propagation properties. Therefore, since the Mn08 alloy has a higher resistance to fatigue crack propagation than that of Mn00 alloy, Mn08 alloy is expected to be in a better position for the structural component applications.
4. Conclusions The homogeneous deformation and dispersion hardening by the means of Mn dispersoids are found to increase the strength of an alloy without sacrificing its ductility and at the same time extends the LCF life. From the results of TEM analysis, it is verified that Mn dispersoids act as a barrier to a dislocation motion and promotes cross slip, with uniform and homogeneous deformation.
1 Di Russo E, Conserva M, Buratti M, Gatto F. A new thermo-mechanical procedure for improving the ductility and toughness of AlZn-Mg-Cu alloys in the transverse directions. Mater Sci Eng 1974;14:23–36. 2 Hahn GT, Rosenfield RA. Metallurgical factors affecting fracture toughness of aluminum alloys. Metall Trans A 1975;6A:653–68. 3 Sanders RE Jr, Starke EA Jr. The effect of grain refinement on the low cycle fatigue behavior of an aluminum-zinc-magnesium(zirconium) alloy. Mater Sci Eng 1977;28:53–68. 4 Wagner JA, Shenoy RN. The effect of copper, chromium, and zirconium on the microstructure and mechanical properties of Al-ZnMg-Cu alloys. Metall Trans A 1991;22A:2809–18. 5 Bryant AJ. The effect of composition upon the quench-sensitivity of some Al-Zn-Mg alloys. J Inst Met 1966;94:94–9. 6 Park DS. Effects of Mn-dispersoid on the deformation mechanism and precipitate free zone on the low cycle fatigue of an Al-Zn-Mg alloy. Ph.D. Thesis, 1991, KAIST. 7 Park DS, Nam SW. On the characteristics of Mn dispersoid in AlZn-Mg alloys. J Mater Sci Lett 1994;13:716–8. 8 Park DS, Nam SW. Effect of manganese on the precipitation in an Al-Zn-Mg alloy. J Mater Sci Lett 1991;10:397–9. 9 Park DS, Nam SW. Effects of manganese dispersoid on the mechanical properties in Al-Zn-Mg alloys. J Mater Sci 1995;30:1313– 20. 10 Sanders TH Jr, Starke EA Jr. The relationship of microstructure to monotonic and cyclic straining of two age hardening aluminum alloys. Metall Trans A 1976;7A:1407–18. 11 Chen F-R, King AH. Determination of the crystallographic directions and planes of features and of the misorientations of crystals with high accuracy and internal estimation of errors. J Elec Micro Technol 1987;6:55–61. 12 Marukawa K. Identification of dislocations and stacking faults by electron microscopy (1). Bull Jpn Inst Met 1984;23:182–91. 13 Ritchie RO, Suresh S. Some considerations on fatigue crack closure at near-threshold stress intensities due to fracture surface morphology. Metall Trans A 1982;13A:937–40. 14 Jo BL, Park DS, Nam SW. Effect of Mn dispersoid on the fatigue crack propagation of Al-Zn-Mg alloys. Metall Mater Trans A 1996;27A:490–3. 15 Jo BL. Effects of PFZ (Precipitate Free Zone) and Mn-dispersoid on the fatigue crack propagation behavior in Al-Zn-Mg-Zr-(Mn) alloy. M.S. Thesis, 1994, KAIST.
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16 Sanders RE Jr., Starke EA Jr.. The effect of ingot processing on the microstructure and fatigue properties of 7475-T6 plates. In: Morris JG, editor. Thermomechanical Processing of Aluminum Alloys. TMS-AIME, 1978:50–73. 17 Grau F-J, Gysler A, Lu¨tjering G. Effect of microstructure on the fatigue crack propagation behavior of high-strength Al-Zn-Mg-Cu alloys. In: Sanders TH Jr., Starke EA Jr., editors. The 4th Inter-
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national Conference on Aluminum Alloys: Their Physical and Mechanical Properties. The Georgia Institute of Technology, 1994:709–16. 18 Hertzberg RW. Deformation and fracture mechanics of engineering materials. New York: John Wiley and Sons, 1986. 19 Fuchs HO, Stephens RI. Metal fatigue in engineering. New York: John Wiley and Sons, 1980.
Enhancement of the fatigue properties in a weldable high-strength Al-Zn-Mg-Mn alloy by means of Mn-dispersoids Duck Hee Lee a, Kyu Cheol Kim b, Dong Seok Park c, Soo Woo Nam
d,*
a
d
Department of Mat. Sci. and Eng., Korea Advanced Institute of Science and Technology, 373-1 Kusong-dong Yusong-gu, Taejon 305-701, South Korea b Andersen Consulting, 17-3 Yoido-dong Youngdeoungpo-ku, Seoul 150-010, South Korea c Defense Quality Assurance Agency, 341 Samjeongja-dong, Changwon 641-160, South Korea Dept. of Mat. Sci. and Eng., Korea Advanced Institute of Science and Technology, 373-1 Kusong-dong Yusong-gu, Taejon 305-701, South Korea Received 10 August 1998; received in revised form 29 October 1998; accepted 19 November 1998
Abstract This present research was undertaken to investigate the effects of Mn dispersoids on low cycle fatigue (LCF) and high cycle fatigue (HCF) in the peak aged Al-Zn-Mg-Mn alloy. The Mn addition in an Al-Zn-Mg-Mn alloy was found to form Mn dispersoids in sizes ranging from 0.05 to 0.5 µm. Mn dispersoids in an Al-Zn-Mg-Mn alloy results in increasing strength without sacrificing ductility by its dispersion hardening effects but at the same time homogeneous deformation. The homogeneous deformation is caused by the cross-slip of dislocations at Mn dispersoids and prevents stress and/or strain concentrations. In this Mn dispersoid containing alloy, both LCF and HCF lives are found to be significantly extended compared with that without any Mn dispersoids. This phenomenon is very interesting and is considered to be due to the simultaneous improvement of strength and ductility by Mn dispersoids. In this investigation, the reason why these excellent properties are observed is discussed in terms of the interaction between Mn dispersoids and the dislocation. 1999 Elsevier Science Ltd. All rights reserved. Keywords: Al-Zn-Mg-Mn alloy; Low cycle fatigue; Fatigue crack growth rate; Mn dispersoids; Cross-slip
1. Introduction The typical commercial 7XXX aluminum alloys are divided into two groups by the characteristic properties, i.e., the one is the weldable aluminum alloys of 7039 and 7017 etc. (UTS: 440–480MPa) which are being produced as thick plate to be used for the structural materials of transportation, aeroplanes, and military vehicles. The other is the non-weldable high strength aluminum alloys of 7075 and 7050 etc. (UTS: 540– 570MPa) which are being produced by extrusion process and these alloys have about 20–25% higher strength compared to that of the weldable alloys. It is well known that a small amount of transient elements such as Zr, Cr, and Mn form various dispersoids which affect the mechanical properties of aluminum
* Corresponding author. Tel.: +82 42 869 3318; fax: +82 42 869 3310; e-mail: [email protected].
alloys.[1–5] With this basic concept, adding more than 0.5wt.% Mn to the weldable Al-Zn-Mg series aluminum alloy, a new weldable high-strength aluminum alloy is developed.[6,7] This new weldable high-strength aluminum alloy has strength as high as that of Al 7075 and at the same time it has good weldability, as good as Al 7039. Therefore, this new alloy can be used to produce many structural parts of transportation, high speed trains, aeroplanes, and military vehicles. Mn dispersoids in a new Al-Zn-Mg-Mn alloy are found to act as the major agents in increasing strength without sacrificing ductility.[8,9] They also do not affect the weldability of this alloy, as well as the kinetics of η⬘ precipitation during the artificial aging process.[8] In general, the fatigue is classified by two regimes, low-cycle fatigue (LCF) and high-cycle fatigue (HCF). These two different fatigue phenomena have been known to be governed by different mechanical properties, i.e., LCF is mainly related with ductility and HCF is strongly affected by strength. Unfortunately, it is gen-
0142-1123/99/$ - see front matter 1999 Elsevier Science Ltd. All rights reserved. PII: S 0 1 4 2 - 1 1 2 3 ( 9 8 ) 0 0 0 8 5 - 1
384
D. H. Lee et al. / International Journal of Fatigue 21 (1999) 383–391
erally accepted that, in metals, increase of strength results in the reduction of ductility. Therefore, it has been reported that commercial aluminum alloys with good LCF life have poor HCF strength and vice versa. In the present study, the role of Mn dispersoids in LCF and HCF behavior of Al-Zn-Mg-Mn alloys is investigated through TEM analysis in combination with fatigue tests.
tensile test machine attached to a tension-compression auto-toggling controlling system. These SEM in-situ fatigue tests were conducted at a strain rate of 1 mm/min and a total strain range of ±2%. Fatigue test specimens were solution treated at 460°C for 2 h, water quenched, and then tested.
3. Results and discussion 2. Experimental procedures
3.1. Microstructure and tensile properties
The chemical compositions of the alloys tested are presented in Table 1. The Al-Zn-Mg-(Mn) alloys were prepared by melting 99.99% pure Al with the appropriate master alloys and casting in an inert atmosphere. After homogenization treatment at 460°C for 24 h which is to form the Mn dispersoids and homogenize the segregation, the billets were extruded at 400°C to be 25 mm × 100 mm plates. Hot rolling of billets to 25 mm × 150 mm plates was also performed at 390°C with a final reduction ratio of 80%. The plates were solution treated at 460°C for 90 min, followed by a water quench. A twostep artificial aging treatment was carried out at 97°C for 8 h and then at 140°C for 10 h to give higher strength and less effective quench sensitivity, compared to the conventional one-step aging treatment. Low-cycle fatigue test specimens which had a gauge length of 6 mm and a gauge diameter of 4.5 mm were machined from the extruded and rolled plates with the stress axis parallel to the extrusion and rolling direction. The fully reversed axial strain-controlled low-cycle fatigue tests were conducted in air using an Instron 1380 with a constant strain rate of 4×10⫺3s⫺1. Fatigue crack growth tests were performed using a 10 ton capacity Instron 8500 dynamic machine at room temperature in air environment, with the procedures outlined in ASTM E647-88. Specimens were cut from the extruded and rolled plates in the LT orientation and machined to be the standard compact tension specimens with chevron notch. All the tests were carried out with 5 Hz triangular frequency under the load ratio of 0.25. Crack length was measured during test using a traveling microscope. For the SEM in-situ observation of the deforming surface morphology, in-situ low-cycle fatigue test specimens were made with a gauge length of 5.5 mm and a diameter of 2.5 mm, and cycled in a JEOL SM TS40
The morphology and distribution of Mn-dispersoids within two different contents of manganese are shown in Fig. 1. In the Mn bearing alloy, with an extruded and rolled Mn08 alloy, spherical- and rod-shaped Mn-dispersoids of the size in range from 0.05–0.5 µm are found to be dispersed uniformly in the matrix. The distribution and morphology of Mn dispersoids in rolled Mn00 and Mn08 alloys are almost consistent with those of extruded ones, respectively, even though the results of rolled alloys are not presented in this paper. It has been experimentally observed that the Mn-dispersoids were formed during a homogenization treatment and no change is
Table 1 Chemical composition of the investigated Al-Zn-Mg-(Mn) alloys(in wt.%) Alloy
Zn
Mg
Mn
Zr
Al
Mn00 Mn08
4.2 4.3
2.6 2.8
0.07 0.8
0.15 0.15
bal. bal.
Fig. 1. TEM micrographs showing the size and distribution of Mn dispersoids in the extruded alloys: (a) Mn00 alloy, (b) Mn08 alloy.
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found in the morphology and size during the aging treatment.[8] Also, according to the observation of Park,[8] Mn-dispersoids in the Al-Zn-Mg-Mn alloy is found not to affect the kinetics of precipitation by the artificial aging process. Mechanical properties of the investigated alloys are tabulated in Table 2. Both the yield and ultimate tensile strength of these Mn bearing alloys are measured to be increased with sufficient amount of elongation. This result implies that the fine Mn dispersoids in the matrix could act as a restriction for any dislocation motion, and consequently increase the strength of the Mn bearing alloy, extruded and rolled. 3.2. Tensile deformation behavior In order to identify the deformation mode of the investigated Al alloys, a dislocation structure was investigated using the specimen strained up to 1% as shown in Fig. 2. The TEM micrographs in Fig. 2 show that there is a trend of slip homogenization with increasing the density of Mn-dispersoids. It appears that this alternation in deformation behavior is due to the Mn dispersoids, since no significant difference in microstructure was found except this distribution of the Mn dispersoids. In Fig. 2(a), Mn00 alloy shows the wide width of intense slip bands consisting of high density of dislocations. However, Mn-dispersoids disperse slip and inhibit the formation of intense slip bands as shown in Fig. 2(b). Therefore, plastic deformation is more homogeneous. In both extruded and rolled conditions, it is believed that the dispersive distribution of Mn-dispersoids in Mn08 alloy promotes homogeneous deformation by converting the deformation mechanism from planar slip, which is well developed in the Mn00 alloy with no Mn dispersoids, to wavy slip. As for the homogeneous deformation in Mn08 alloy, the deformation mode could be mainly dominated by cross slip.[10] To identify the interaction behavior between Mn dispersoid and dislocation, TEM analysis was introduced with the specimen imposed small amount of plastic deformation which will be in detail discussed in the ‘TEM analysis’ section (Fig. 7). From the above TEM observation and previous works,[6,9] the role of Mn dispersoids in the Mn08 alloy is very clearly seen in that they act as a strengthener to block the motion of Table 2 Tensile properties Alloy
Condition
YS (MPa)
UTS (MPa) El. (%)
Mn00
Extruded Rolled Extruded Rolled
420 375 484 408
463 445 545 468
Mn08
15.2 13.0 12.3 14.3
Fig. 2. Comparison of the distribution of slip band after 1% tensile strain in the extruded alloys: (a) Mn00 alloy, B=[111], g=⬍220⬎; (b) Mn08 alloy, B=[111], g=⬍220⬎.
dislocations, and at the same time they distribute the slip of dislocations uniformly so as to enhance ductility. 3.3. Low cycle fatigue A Coffin-Manson plot is shown in Fig. 3. In this figure, we can easily see that the Mn bearing alloy, extruded and rolled, has a longer LCF life than no Mn bearing alloy in both extruded and rolled conditions, even though the strength of the Mn08 alloy is higher than that of the Mn00 alloy. Usually, a metal having higher elongation or ductility with lower strength has a longer LCF life than a low-elongation metal with high strength because the LCF test mode is strain-controlled. However, the Mn08 alloy shows sufficient ductility with higher strength but also has longer LCF life, compared with the Mn00 alloy. Cyclic properties of the investigated alloys are tabulated in Table 3. In order to explain the above mentioned results of the LCF life, the dislocation structure was investigated using the specimens fatigued up to 10% of the LCF life, as shown in Fig. 4. The result in Fig. 4 is similar to that observed in a previous section which was obtained
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Fig. 3.
Coffin-Manson plot of each alloy.
Table 3 Cyclic properties of low cycle fatigue specimens (Nfα·⌬⑀p = C, C = constant) Alloy
Condition
C
⫺a
Mn00
Extruded Rolled Extruded Rolled
0.188 0.239 0.263 0.402
0.667 0.670 0.688 0.728
Mn08
Fig. 4. TEM micrographs showing the dislocation arrangement formed after cyclic deformation in the rolled alloys: (a) Mn00 alloy, (b) Mn08 alloy.
throughout tensile deformation. When the inhomogeneous deformation occurs by planar slip in the Mn00 alloy, as shown in Fig. 4(a), the stress concentration around the pile-up dislocations at the end of slip band may occur and interact with grain boundaries. This local stress concentration may cause early fatigue crack initiation at grain boundaries which would then decrease LCF life. On the other hand, since the Mn dispersoids make the deformation homogeneous, as shown in Fig. 4(b), fatigue crack initiation due to the local stress concentration at grain boundaries is suppressed in the Mn bearing Mn08 alloy. Therefore, it appears that homogeneous deformation by means of the Mn dispersoids will lead to a longer LCF life. These phenomena will be in detail discussed in the following section. 3.4. SEM in-situ fatigue test The results of in-situ fatigue tests of the Mn00 alloy show typically coarse slip bands formed at 45° angles to the stress axis and often concentrated at grain boundary areas, as shown in Fig. 5. The slip band traces impinged at the grain boundaries may build up local stresses that can cause fatigue crack initiation. These
cracks are continuously extended and coalesce together during further cycling. In Fig. 5(c), the enlarged grain boundary cracks after 1100 cycles are shown. The localized slip bands in this alloy without Mn were found to be caused by a planar dislocation movement (Fig. 4(a)). The corresponding results of in-situ fatigue testing of the Mn08 alloy are shown in Fig. 6. This alloy containing Mn-dispersoids does not show any well-developed slip bands throughout its fatigue life. Thus it is believed that this kind of deformation mode is caused by the Mndispersoids. They omit coarse slip band formation and cause homogeneous deformation. As the number of fatigue cycles increases, small cracks are seen all over the specimen’s surface. These cracks are distributed, not at grain boundaries, but inside the grains. This indicates that they were neither caused nor affected by the concentrated slip bands close to the grain boundaries. This homogeneous deformation of the Mn08 alloy makes crack initiation difficult and extends its fatigue life.
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Fig. 5. SEM micrographs from in-situ fatigue tests of the Mn00 alloy at ⌬⑀t=±2%. Arrows indicate fatigue stress axis: (a) after 120 cycles, (b) after 790 cycles (c) after 1100 cycles.
3.5. TEM analysis To verify the cause of homogeneous deformation mode by Mn-dispersoids during fatigue testing, TEM analysis on the dislocation structure was conducted for the Mn08 alloy. The observations were made after only one cycle to avoid a very complicated dislocation structure due to multiplication and tangling. The TEM micrographs and corresponding schematic diagrams, which show the 3-D configuration of dislocations, are shown in Fig. 7. In order to analyze the dislocation, before and after its interaction with the Mn-dispersoids during fatigue testing, trace analysis was applied.[11,12] Obtaining the Burgers vector, the usual g·b criterion was used, and for the line direction determination, 3 beam directions were chosen to get the diffraction pattern. Using the relationship between the projected dislocation and that of the real, the unknown value of the real dislocation direction can be obtained from the known projected one.[11] The slip plane of the dislocations, containing both Burgers vector and dislocation line vector, can be easily obtained by the cross product of the two
vectors. Calculated Burgers vector, line vector and slip plane are listed in Table 4. In Fig. 7, dislocations are found to be moving on a (111) slip plane (a, b and c) until they are blocked by the non-shearing Mn-dispersoids and the dislocations tend to cross-slip to (111) slip plane (d and e). This indicates that dislocations crossslip from primary slip plane, (111), to the other active slip plane, (111), by the Mn-dispersoids without any dislocation pile-up at the Mn-dispersoid. Such a phenomenon of blocking and cross-slip of dislocations corresponds with the SEM in-situ test results in that the slip steps are very fine and evenly distributed. This suggests slip homogenization through the cross-slip over the nonshearable obstacles. Therefore, it is believed that the dispersive distribution of the coarser Mn-dispersoids in a Mn08 alloy promotes homogeneous deformation by converting the dislocation movement from planar slip to a wavy slip mode through cross-slip. This results in increasing strength without sacrificing ductility, as well as retarding fatigue crack initiation and propagation. All of this is generated by the inhomogeneous deformation within the grain boundary.
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Fig. 6. SEM micrographs from in-situ fatigue tests of the Mn08 alloy at ⌬⑀t=±2%. Arrows indicate fatigue stress axis: (a) after 170 cycles, (b) after 800 cycles (c) after 1200 cycles.
Fig. 7. TEM micrograph showing the dislocation structure of the Mn08 alloy after 1 cycle at ⌬⑀t=±1.7%. (B=[112], g=[220]; Arrows indicate Mn-dispersoids).
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Table 4 Dislocation characters and slip planes in the Mn08 alloy Dislocation b a b c d e
1/2[0 1/2[0 1/2[0 1/2[0 1/2[0
1¯ 1¯ 1¯ 1¯ 1¯
1] 1] 1] 1] 1]
x
Slip plane
Slip system
[61 50 11] [152 117 35] [9 15 6¯] [23 8 15] [9 16 7]
(1¯ (1¯ (1¯ (1 (1
primary primary primary secondary secondary
1 1 1 1 1
1) 1) 1) 1) 1)
3.6. High-cycle fatigue For the high-cycle fatigue test, the fatigue crack growth rate has been measured and plotted in terms of ⌬K, as shown in Fig. 8. In Fig. 8, the threshold stress intensity (⌬Kth) values were obtained from the plot of da/dN vs. ⌬K within the region I suggested by Ritchie et al.[13] The value of ⌬Kth is considered to strongly depend on alloy composition, ingot process, thermomechanical treatment, test condition and so on.[14–17] The fatigue crack growth tests at several different conditions including that presented in this article were conducted using the Mn00 and Mn08 alloys. Under all test conditions, it was found that threshold values occurred at about 10⫺7 m/cycle.[14,15] Therefore, even though the ⌬Kth values in the present study are thought to be somewhat higher than those of commercial Al alloys, this finding seems to be characteristic of these tested alloys. In addition, from the literature survey,[17, 18 (pp. 519– 618), 19 (pp. 56–104)] it can be seen that as the threshold value of a material is higher, the threshold value
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generally occurs at higher growth rate (higher da/dN) although this event is not always true. The Mn08 alloy is found to have a larger threshold stress intensity than the Mn00 alloy in both extruded and rolled conditions. Also, at the same stress intensity which was over the threshold stress intensity, the Mn08 alloy has a much slower crack growth rate compared to that of the Mn00 alloy and thus, the fatigue crack growth resistance of the Mn08 alloys is superior to those of the Mn00 alloys. These results are thought to be due to Mn dispersoids, since the major difference between Mn bearing and none Mn containing alloys is whether there are Mn dispersoids or not. The fatigue crack growth properties of these investigated aluminum alloys are summarized in Table 5. On the other hand, the difference in fatigue crack growth properties between the extruded and rolled alloys is thought to be due to the aging treatments which are considered to be a major affecting factor and grain structures, because these are the main differences for two kinds of alloys with the different treatments (extrusion and rolling). The different aging treatment was conducted on extruded specimens for the fatigue crack growth test (Fig. 8). Extruded alloys (Mn00 and Mn08 alloy) have the partially recrystallized grains within the elongated fibrous grains along the extrusion direction. As the content of Mn increases, the recrystallized grain size somewhat decreases. The grain sizes of Mn00 and Mn08 alloys are about 13 µm and 8 µm, respectively. However, the rolled alloys have the coarser pancake shaped grain structure (⬎30 µm) than those of the extruded alloys with same Mn composition. To identify the above mentioned results, the fracture surfaces in extruded alloys were examined by SEM. Fig. 9 shows the differences in the morphology of the fracture surfaces. The extruded Mn08 alloy has a much more rough fracture surface. Therefore, in the Mn08 alloy, it is supposed that as the fatigue crack advances, the Mn dispersoids may cause crack deflections and make the crack path more tortuous by acting as an obstacle to the advance of a fatigue crack. These results illustrate that Mn dispersoids can increase many more deflections in the fatigue crack path than the inhomogeneous deformation, causing the fatigue crack propagation curve to shift from the left to right, as illustrated in Fig. 8. Table 5 Fatigue crack propagation properties of the investigated Al alloys (da/dN = A(⌬K)m)
Fig. 8. Influence of the Mn content on the fatigue crack propagation behaviors. (Extruded alloys: natural aged at R.T./96h → 100°C/10 min +160°C/180 min; YS(MPa)/UTS(MPa)/El.(%) =258/344/20 (Mn00 alloy), 360/428/16.5 (Mn08 alloy)).
Alloy
Condition
⌬Kth (MPa·m1/2)
A m (×10⫺11)
Mn00
Extruded Rolled Extruded Rolled
8.22 15.46 10.18 18.38
1.70 1.75 1.92 0.06
Mn08
3.74 3.59 3.68 4.32
390
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The Mn dispersoids in the Mn08 alloy increase the tortuosity of the crack path and improve the fatigue crack growth properties, through a crack deflection caused by the finely dispersed Mn dispersoids. Considering the results of the LCF and HCF tests in the extruded and rolled Mn bearing alloy, it can be seen that finely dispersed Mn dispersoids can improve not only the LCF life but also resistance to fatigue crack growth.
Acknowledgements This work was supported by MOST in the regime of the HAN (NM05630) project. Thanks are given for its financial support.
References
Fig. 9. SEM micrographs showing the fatigue crack growth path and fracture surface in the extruded alloys: (a) Mn00 alloy (b) Mn08 alloy.
The structural components have notches of different acuity. In the view point of notch sensitivity of a material, the sharp notch may cause earlier fatigue short crack initiation at the notch tip. However, once the fatigue short crack is, to some extent, extended and not affected by stress field related with notch, the sharpness of notch may have less effect on crack propagation properties. Therefore, since the Mn08 alloy has a higher resistance to fatigue crack propagation than that of Mn00 alloy, Mn08 alloy is expected to be in a better position for the structural component applications.
4. Conclusions The homogeneous deformation and dispersion hardening by the means of Mn dispersoids are found to increase the strength of an alloy without sacrificing its ductility and at the same time extends the LCF life. From the results of TEM analysis, it is verified that Mn dispersoids act as a barrier to a dislocation motion and promotes cross slip, with uniform and homogeneous deformation.
1 Di Russo E, Conserva M, Buratti M, Gatto F. A new thermo-mechanical procedure for improving the ductility and toughness of AlZn-Mg-Cu alloys in the transverse directions. Mater Sci Eng 1974;14:23–36. 2 Hahn GT, Rosenfield RA. Metallurgical factors affecting fracture toughness of aluminum alloys. Metall Trans A 1975;6A:653–68. 3 Sanders RE Jr, Starke EA Jr. The effect of grain refinement on the low cycle fatigue behavior of an aluminum-zinc-magnesium(zirconium) alloy. Mater Sci Eng 1977;28:53–68. 4 Wagner JA, Shenoy RN. The effect of copper, chromium, and zirconium on the microstructure and mechanical properties of Al-ZnMg-Cu alloys. Metall Trans A 1991;22A:2809–18. 5 Bryant AJ. The effect of composition upon the quench-sensitivity of some Al-Zn-Mg alloys. J Inst Met 1966;94:94–9. 6 Park DS. Effects of Mn-dispersoid on the deformation mechanism and precipitate free zone on the low cycle fatigue of an Al-Zn-Mg alloy. Ph.D. Thesis, 1991, KAIST. 7 Park DS, Nam SW. On the characteristics of Mn dispersoid in AlZn-Mg alloys. J Mater Sci Lett 1994;13:716–8. 8 Park DS, Nam SW. Effect of manganese on the precipitation in an Al-Zn-Mg alloy. J Mater Sci Lett 1991;10:397–9. 9 Park DS, Nam SW. Effects of manganese dispersoid on the mechanical properties in Al-Zn-Mg alloys. J Mater Sci 1995;30:1313– 20. 10 Sanders TH Jr, Starke EA Jr. The relationship of microstructure to monotonic and cyclic straining of two age hardening aluminum alloys. Metall Trans A 1976;7A:1407–18. 11 Chen F-R, King AH. Determination of the crystallographic directions and planes of features and of the misorientations of crystals with high accuracy and internal estimation of errors. J Elec Micro Technol 1987;6:55–61. 12 Marukawa K. Identification of dislocations and stacking faults by electron microscopy (1). Bull Jpn Inst Met 1984;23:182–91. 13 Ritchie RO, Suresh S. Some considerations on fatigue crack closure at near-threshold stress intensities due to fracture surface morphology. Metall Trans A 1982;13A:937–40. 14 Jo BL, Park DS, Nam SW. Effect of Mn dispersoid on the fatigue crack propagation of Al-Zn-Mg alloys. Metall Mater Trans A 1996;27A:490–3. 15 Jo BL. Effects of PFZ (Precipitate Free Zone) and Mn-dispersoid on the fatigue crack propagation behavior in Al-Zn-Mg-Zr-(Mn) alloy. M.S. Thesis, 1994, KAIST.
D. H. Lee et al. / International Journal of Fatigue 21 (1999) 383–391
16 Sanders RE Jr., Starke EA Jr.. The effect of ingot processing on the microstructure and fatigue properties of 7475-T6 plates. In: Morris JG, editor. Thermomechanical Processing of Aluminum Alloys. TMS-AIME, 1978:50–73. 17 Grau F-J, Gysler A, Lu¨tjering G. Effect of microstructure on the fatigue crack propagation behavior of high-strength Al-Zn-Mg-Cu alloys. In: Sanders TH Jr., Starke EA Jr., editors. The 4th Inter-
391
national Conference on Aluminum Alloys: Their Physical and Mechanical Properties. The Georgia Institute of Technology, 1994:709–16. 18 Hertzberg RW. Deformation and fracture mechanics of engineering materials. New York: John Wiley and Sons, 1986. 19 Fuchs HO, Stephens RI. Metal fatigue in engineering. New York: John Wiley and Sons, 1980.