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The effect of microstructure and environment on fatigue crack growth in 7049 aluminium alloy at negative stress ratios
International Journal of Fatigue 25 (2003) 1209–1216 www.elsevier.com/locate/ijfatigue
The effect of microstructure and environment on fatigue crack growth in 7049 aluminium alloy at negative stress ratios M. da Fonte a,∗, F. Romeiro b, M. de Freitas c, S.E. Stanzl-Tschegg d, E.K. Tschegg e, A.K. Vasude´van f a
Nautical School, 2780-572 Pac¸o de Arcos, Portugal b Instituto Polite´cnico de Leiria, ESTG, Portugal c Instituto Superior Te´cnico, Technical University of Lisbon, Lisbon, Portugal d Institute of Met and Physics, University of Agricultural Sciences, Turkenschanzstrasse 18, A-1180 Vienna, Austria e Institute of Applied and Technical Physics, Technical University of Vienna, Karlsplatz 13, A-1040 Vienna, Austria f Office of Naval Research Laboratory, Arlington, VA 22217, USA
Abstract The influence of environment and microstructure on fatigue crack growth has been investigated on a high strength 7049 aluminium alloy. This aluminium alloy was artificially aged to underaged (UA) and overaged (OA) microstructures. The heat treatment procedure was performed in order to obtain an UA and OA microstructure having the same yield strength properties, but differing in the mode of slip deformation: the UA alloy deforms by planar slip and that of the OA alloy by wavy slip. The crack growth measurements were performed in MT specimens at constant load ratios for R = 0, ⫺1, ⫺2, ⫺3 near-threshold and Paris regime in ambient air and vacuum conditions. Crack closure loads were measured in order to determine the Popen for each R ratio. Micromechanisms of near-threshold crack growth are briefly discussed for several concurrent processes involving environmentally assisted cracking with intrinsic microstructural effects. The results showed that the presence of humid air leads to a larger reduction in ⌬Kth for both the ageing conditions, but the UA specimens were superior probably because of crack branching. The role of environmental effect and microstructures near-threshold regime seems to be more significant than any mechanical contributions to the crack closure, such as plasticity, roughness, oxide, etc. 2003 Elsevier Ltd. All rights reserved. Keywords: Microstructure; 7049 and 7075 alloys; Slip mode; Near-threshold fatigue; Fatigue crack growth; Load ratio; Crack closure
1. Introduction The majority of investigations to date on the fatigue crack growth resistance of high strength aluminium alloys has focussed on crack advance in the so-called Paris regime (da / dNⱖ10⫺9 m / cycle). However, the behaviour in the near-threshold fatigue regime, which accounts for a significant portion of the crack propagation life, should deserve more attention when compared with the numerous studies concerning with crack closure effects. Designers and the industry in general need to obtain from the researcher intrinsic parameters such as the stress intensity factor range ⌬K and the
∗
Corresponding author. Fax: +351-21-4429-546. E-mail address: [email protected] (M. da Fonte).
0142-1123/$ - see front matter 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0142-1123(03)00150-6
maximum stress intensity factor Kmax to characterise the fatigue crack growth behaviour of materials under service loading conditions, microstructures and environment. They cannot be dependent on extrinsic testing parameters, such as crack closure and environmental effects, which are always very difficult to obtain. Synergetic effects of environment and loading make the understanding of the underlying mechanisms between microstructure and environment difficult. The microstructure and the environment play an important role in the fatigue crack growth resistance of Al alloys and have been investigated extensively during the last two decades [1–9]. Kirby and Beevers [10], on the other hand, clearly demonstrated that even the seemingly innocuous environment of laboratory air can lead to a marked increase in crack propagation rates, as compared to vac-
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uum, in the near-threshold fatigue regime of 7xxx series aluminium alloys. Lin and Stark [11] showed that microstructure–environment interactions at low fatigue crack growth rates may be completely different from those at higher fatigue crack growth rate levels. The concept of fatigue crack closure, introduced in the 1960s by Elber [12], has been used to explain a wide range of positive stress ratio crack propagation results (Rⱖ0), but less attention has been given to fatigue loading for negative stress ratios (R ⬍ 0). The crack closure concept has served to explain most of the related phenomena with decreasing crack growth rates, due to plastic deformation at crack wake roughness of fatigue surfaces, oxide, etc. Crack branching may likewise lead to reduced crack growth rates [13]. The measure of crack closure value has been until now a problem of difficult resolution and Vasude´ van and Sadananda [14–17] have pointed out pertinent problems concerning the subject, having reanalysed the crack closure concepts using dislocation theory. In the absence of closure, the crack tip driving force is related to the applied stress intensity, ⌬Kapp = Kmax⫺Kmin, and when closure is induced, the driving force is decreased to a smaller effective stress amplitude, ⌬Keff = Kmax⫺Kcl. They stated that for a crack propagation to occur, two driving forces, ⌬K and Kmax, are needed simultaneously, instead one. These two driving forces are intrinsic parameters of a material and they are valid for short as well as long cracks. The apparent differences in the behaviour of short or long cracks are related with residual stress. They extended these results in terms of intrinsic behaviour in contrast to the crack closure concepts. They concluded that crack closure has a small effect on fatigue crack growth behaviour, when compared with microstructural and environmental effects. Freitas and colleagues [18] have measured crack closure effects for negative stress ratios in a medium carbon steel DIN Ck 45 for a wide range of stress ratios from R = 0.7 to R = ⫺3. For a given negative R ratio, the Popen depends strongly on the maximum load Pmax. The ratio P open / P max becomes more negative for higher negative stress ratios but only for high applied maximum load levels Pmax. An example of the less importance of the crack closure phenomena when compared with the environment effect (in ambient air and vacuum) is shown in Fig. 1. If the closure due to only wedging action is assumed in the wake of a crack without corrosion influences, then the environmental effect in Ti–6V–4Al alloy data [19] should lie to the right of the vacuum curve. The air data in Fig. 1 lying to the left of vacuum indicate the important role of environment damage ahead of crack tip in the Ti–Al alloy. This is not experimentally observed when the environment results are compared to the results in a good vacuum. The arrow to the left of vacuum curve (square points) indicates the environ-
Fig. 1. Influence of environment on fatigue crack growth in Ti–6V– 4Al alloy [19].
ment’s contribution to fatigue in ambient air. The arrow to the right (dashed curve) of vacuum curve indicates the trend of the moist air curve, if the closure was significant. It has been experimentally observed that da / dN nearthreshold to higher growth rates are nearly independent of R ratios [19,20] in a high vacuum, which leads to a clear conclusion that plasticity does not induce closure as has been postulated three decades ago. Thus, one can infer that the role of environment near-threshold regime is strongly more significant than any mechanical contributions such plasticity, roughness, oxide, etc. Keeping in mind these features, one may conclude that ⌬Kth decreasing with R is an intrinsic fatigue property of the material for that environment. This work presents a systematic fatigue crack growth data from the near-threshold region to the Paris region in a high strength Al 7049 aluminium alloy artificially aged to underaged (UA) and overaged (OA) conditions. The objective was to obtain two microstructures having approximately the same crystallographic texture and yield strength properties, but different in the slip mode. Both macroscopic and microscopic responses are compared at different R ratios and different environments, in air and vacuum at load ratios R between R = 0 and R = ⫺3. 2. Material and experimental procedure Testing material was a high strength 7049 aluminium alloy. Specimens were cut from a cold rolled sheet of a 7049 aluminium alloy with 10 mm thickness. After solution heat treatment (470 °C, 30 min) of the cold rolled plates, one part of the material was underaged at 117 °C for 90 min and the other part was overaged for 8 h at 107 °C plus for 65 h at 163 °C. The chemical composition is
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Table 1 Composition (wt.%) of Al 7049 alloy Zn
Mg Cu
Cr
7.1
2.8
0.06 0.3
1.7
Fe
Si
Mn
Ti
Ga
0.1
0.06
0.05 0.01
Zr
Al
0.1
Bal.
listed in Table 1. The UA and OA heat treatments were designed to produce approximately the same yield strength, in order to eliminate differences in fatigue crack growth behaviour due to the possible strength effects. The yield strengths of both the microstructures are ⬇440 MPa and the ultimate tensile strength of the UA material is 16% higher and its ductility is 100% higher than that of the OA alloy (Table 2). The grain size is up to several hundred micrometres in longitudinal and a few micrometres in transverse direction. Transmission electron micrographs of the different ageing conditions are shown in Fig. 2 [4] for AL 7075 aluminium alloy closed to this Al 7049. The UA material exhibits extremely fine grain precipitate (GP) zones and h⬘ intermetallic precipitates, whereas the OA structure contains predominantly coarse h as well as h⬘ precipitates. h precipitates may be seen at grain boundaries. Tests were carried out in two different servo-hydraulic testing machines, one for testing in vacuum (Schenck, 20 kN) and another one for testing in ambient air (INSTRON, 100 kN). The crack growth measurements were performed in MT specimens (150 × 40 × 5 mm) at constant load ratios ranging from R = 0 to R = ⫺3. The test frequency was 15 Hz. Crack growth rates and ⌬K were obtained according to the standard test method for measurement of fatigue crack growth rates (ASTM E 647—1999). The central notch was 10 mm long and was made by electrical-discharge machining with h = 0.25 mm. A fatigue precracking was conducted till the crack length was approximately 1 mm for each side of the notch. The thickness of the specimen was chosen due to practical reasons related to the machine load capacity of servo-hydraulic Schenck machine (20 kN). However, we did not find any buckling problems under compressive load conditions. Crack growth measurements were made by either an optical microscope with a magnification of 50×, associated with a stroboscopic illumination at the frequency of testing or through an automatic measure-
Fig. 2. Microstructure of Al 7075 alloy [4]: (a) UA Al 7075 alloy, (b) OA Al 7075 alloy.
ment technique using the compliance method associated with load–displacement curves previously calibrated using the data from optical measurements. Special grips were designed and constructed in order to put the MT specimens inside the vacuum chamber, as seen in Fig. 3. The applied maximum loads for the OA and the UA aluminium alloy, testing in air and vacuum, were 5–6 and 7–8 kN, respectively.
Table 2 Mechanical properties of Al 7049 alloy Alloy/Temper
Yield strength Rp
7049-UA 7049-OA
445 441
0.2
(MPa) UTS (MPa) 578 497
1
Elongation (%)
Reduction area (%)
KIC (MPa m2)
17.2 8
19.1 23
– 32.0
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3. Results and discussion Crack growth data were calculated according to ASTM E 647—1999, which states: ⌬K ⫽ Kmax⫺Kmin for R ⬎ 0⌬K ⫽ Kmax for R
(1)
ⱕ0 Therefore: ⌬P ⫽ Pmax⫺Pmin for R ⬎ 0⌬P ⫽ Pmax for Rⱕ0 Results can be analysed as plots of da / dN as a function of ⌬K = Kmax. The calculation of ⌬K for MT specimens is obtained by the following expression: ⌬K ⫽
⌬P B
冪2Wsec 2 pa
pa
(2)
where a = 2a / W. The basic principle of crack closure concept is that a crack grows only if the crack is completely open. Therefore, the stress intensity factor range ⌬K should be reduced to an effective ⌬Keff. The Paris law can be written as Elber law [12]: da / dN ⫽ C(⌬Keff)m
(3)
This is a unified crack growth law for different stress ratios R, where ⌬Keff ⫽ Kmax⫺Kop Fig. 3. Chamber of vacuum with the MT specimen mounted.
Crack closure loads were measured with the compliance technique at test frequency, using a data acquisition system for a wide range of stress ratios 0ⱖRⱖ ⫺3. The centre notched MT specimens were machined after the heat treatments such that crack propagation was parallel to the rolling direction. The grips for the MT specimens used in the chamber of vacuum were specially designed and constructed such that the alignment was perfect and no bending could take place. The experiments were performed in ambient air at 20 °C and at approximately 50% relative humidity. The vacuum was approximately 3 × 10⫺3 Pa. Tests were carried out on a servo-hydraulic testing machine under constant amplitude load control with a 15 Hz sine wave. Threshold ⌬Kth values were defined as those stress intensities, where the maximum crack propagation rates did not exceed 10⫺11 m/cycle. Automatic data acquisition of load (kN) and centreline displacement (mm) was done at regular intervals, such that one recorded results in 2 × 2000 pairs of data points load–displacement.
(4)
To use this unified crack growth law, Kopen or Popen should be determined from experimental data for the different stress ratios, R and Kmax. The use of crack closure (open) concept allows a unique closure free curve for the crack growth rates at different stress ratios to be determined. The measured fatigue crack propagation rates (da / dN vs. Kmax) are shown in Figs. 4 and 5 for loading at
Fig. 4. da / dN vs. Kmax at R = 0, ⫺1, ⫺2, ⫺3 in air.
M. da Fonte et al. / International Journal of Fatigue 25 (2003) 1209–1216
Fig. 5.
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da / dN vs. Kmax at R = 0, ⫺1, ⫺2, ⫺3 in vacuum.
R = 0, ⫺1, ⫺2, ⫺3, in vacuum and in air. The calculation of ⌬K was performed according to ASTM E 647—1999: at negative stress ratios, ⌬K = Kmax. Fig. 4 shows the curves da / dN vs. Kmax for OA and UA aluminium alloy in ambient air conditions near threshold and Paris region. Fig. 5 shows the da / dN vs. Kmax for OA and UA aluminium alloy in vacuum conditions near threshold and Paris region. When comparing the two graphs, Figs. 4 and 5, we can see how significant the effect of microstructure and environment is on crack growth rates near-threshold regime. In vacuum, for both microstructures, UA and OA, thresholds are clearly higher than those in ambient air. The K thresholds and crack growth data are clearly different for OA conditions, and it is justified by the highest level of Pop for the UA alloy, as is shown in Fig. 6a–d. If we adopt for OA a compliance offset of 1%, Pmax / Pop vary from 0.25 for R = ⫺3 up to 0.5 for R = 0, while for UA, Pmax / Pop vary from 0.32 for R = ⫺3 up to 0.55 for R = 0. These figures also show that the crack opening level is similar during crack growth at stress ratios lower than R = ⫺1. It is to be noted that the opening load was determined using the compliance offset method described at ASTM E 647—1999, without filtering the data, which gives a scatter of the opening load level. Despite this scatter, the data show clearly that opening loads are higher for the stress ratio of R = 0 than for negative stress ratio for both alloys and that opening load remains approximately constant for stress ratio lower than R = ⫺1. This confirms the results of Figs. 4 and 5 when crack growth rate is constant from R = ⫺1 to R = ⫺3. It also must be remarked that at stress ratios R = ⫺2 to R = ⫺3, the ASTM practice may not be reasonable, since it states that the compliance that corresponds to the fully open crack confirmation is determined at 25% of cyclic load range on the unloading curve. In
Fig. 6. Determination of opening load (ASTM E 6∗47—1999 method) for the OA (a, c) and the UA (b, d) conditions, at R = 0, ⫺1, ⫺2, ⫺3.
these tests and using this statement, the crack may already be closed. There are no significant differences on crack growth rate for high negative R ratios (R = ⫺2, ⫺3,⫺4). The points are parallel and very close to the curves obtained on Paris region. The strong difference is for R = ⫺1 and R = 0. The measurement of P / P max vs. compliance offset shows these differences clearly for the OA. The influence of environment, i.e. humid air, is plotted for both alloys for a stress ratio of R = ⫺1, ⫺2, ⫺3 in Fig. 4. Two effects are obvious. Firstly, the thresholds are lower in the OA than in the UA conditions and the crack growth rates are higher for the same Kmax; and secondly, the slope of the Paris line is reduced in comparison to UA and the vacuum curves. Table 3 shows the effect of microstructure and environment on the slope m of the Paris law. Table 3 Slope m of Paris law at negative R ratios, in air and vacuum R ratios
da / dN = C (⌬K)m (UA) m
⫺3, ⫺2, ⫺1, 0
Air 7.2
(OA) m Vacuum 8.9
Air 5.1
Vacuum 6.2
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Fig. 7. Fracture surfaces after fatigue loading in the threshold regime. (a) OA in air, R = ⫺3; (b) OA in vacuum, R = ⫺3.
Fig. 8. Fracture surfaces after fatigue loading in the threshold regime. (a) UA in air, R = ⫺1; (b) UA in vacuum, R = ⫺1.
Figs. 7–9 show the fracture surfaces of the OA and UA 7049 aluminium alloy after loading at negative R ratios on the threshold regime in air and vacuum. The different negative R ratios did not result in different fracture surface appearance. This result corresponds to that of the identical crack growth curves. The reproduced figures therefore are considered to represent the structures of near-threshold tested specimens at all tested negative R ratios. In Fig. 7, the fracturing features of the OA alloy after loading at R = ⫺3 in air (Fig. 7a) and vacuum (Fig. 7b) are compared. Both fractures appear smooth, ductile and flat. In the UA condition, however, a more brittle fracture, with crystallographic features, is observed after loading in air (Fig. 8a). These features were eventually observed likewise in specimens, which have been loaded in vacuum, but there mostly smooth and ductile features prevail (Fig. 8b). In addition, secondary cracks, which point to crack branching, are recognised (Fig. 9a). Crack branching was indeed observed during the measurements of the UA alloy in air and is considered as an important mechanism for crack growth retardation.
Fig. 9b was taken with a tilt angle of 35° in the SEM and thus makes the high fracture surface roughness visible. The higher fracture surface roughness of the UA specimens could be another reason for the better fatigue crack growth properties of this alloy. The reported results show that the environment and the microstructure strongly influence the fatigue crack growth resistance of high strength aluminium alloys as is known from the literature. The microstructure and the applied ⌬K primarily control the slip mode for crack growth, which can be significantly changed by the presence of the environment. The apparent differences in fatigue crack growth resistance of the two ageing conditions are ascribed to a complex interaction between two mechanisms: the embrittling effect of humid air resulting in conventional corrosion fatigue process, the role of microstructure and slip mode in inducing crack deflection and crack closure arising from a combination of environment and microstructural contributions.Corrosion fatigue processes, i.e. active path corrosion and hydrogen embrittlement [11,21,22], tend to accelerate crack growth. Although such mechanisms are not clearly understood for alu-
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slope m significantly different for both microstructures (OA and UA) and different environment conditions, in air and vacuum (see Table 3).
4. Conclusions
Fig. 9. Fracture surfaces after fatigue loading in the threshold regime. (a) UA in air, R = ⫺1; (b) UA in vacuum, R = ⫺3 (tilt 35°).
minium alloys, it is generally accepted that hydrogen embrittlement induced by water vapour with freshly formed fracture surface is responsible for the resulting enhancement of crack propagation rates in comparison with an active path dissolution mechanism. The results of this study show that the presence of humid air leads to a significant reduction of Kmax, Kth for both ageing conditions (OA, UA) when compared with the testing vacuum conditions (3 × 10⫺3 Pa). There is no evidence for the oxide induced closure. The apparent differences in fatigue crack growth resistance of the two microstructures are ascribed to a complex interaction between several mechanisms such as change of fracture mode due to different microstructure and slip mode, as well as hydrogen embrittlement. The different microstructures manifested in terms of their different deformation slip mode of planar vs. wavy indicate that the resistance to crack growth in a planar slip alloy is significantly improved in comparison to that of the wavy slip alloy, with crack branching playing an important role in the underaged alloy. Another result concerns the C and m parameters of the Paris law, which are traditionally considered as characteristic material properties. The results showed a
The fatigue crack growth threshold behaviour of an Al 7049 alloy was studied by using a material with identical chemical composition and yield strength, but different microstructure. Experiments at negative R ratios and different environment (ambient air and vacuum) showed that Kmax thresholds depend on the different microstructure and the associated deformation mechanisms. These are homogeneous slip in the OA condition and localised slip with crystallographic cracking and with a tendency to crack branching in the UA material. The second most important influence comes from the environment. It shows that the Kmax threshold is reduced by approximately 50%, probably by hydrogen embrittlement. The underaged and overaged microstructures of this alloy show significant differences in the fatigue crack growth resistance at near-threshold crack growth rates. There are no significant differences on crack growth rates for higher negative R ratios either in ambient air or in vacuum. The significant difference appears between R = 0 and R = ⫺1, which is demonstrated by the crack opening load shown in the P / Pmax compliance offset. The parameters C and m of the Paris law, traditionally considered as a specific property of materials, significantly differ for each microstructure and environment.
References [1] Hornbogen E, Starke EA. Theory assisted design of high strength low alloy aluminium (overview). Acta Metall Mater 1991;4:1–16. [2] Ming Gao, Pao PS, Wei RP. Chemical and metallurgical aspects of environmentally assisted fatigue crack growth in Al 7075-T651 alloy. Metall Trans 1988;19A:1739. [3] Lafarie-Frenot MC, Gasc C. The influence of age-hardening of fatigue crack propagation behaviour in aluminium alloy in vacuum. Fatigue Fract Eng Mater Struct 1983;6:329. [4] Suresh S, Vasude´ van AK, Bretz PE. Mechanism of slow fatigue crack growth in high strength aluminum alloys. Metall Trans A 1984;15A:369–79. [5] Davidson D, Lankford J. The effects of aluminum alloy microstructure on fatigue crack growth. Metall Sci Eng 1985;74:189–99. [6] Vasude´ van AK, Suresh S. Influence of corrosion deposits on near threshold fatigue crack growth behaviour in 2xxx and 7xxx series aluminum alloys. Metall Trans A 1982;13A:2271–80. [7] Petit J, Zeghloul A. In: Backlund J, Blom A, Beevers CJ, editors. Fatigue thresholds. Warley, UK: EMAS Ltd; 1982. p. 56-3. [8] Tschegg-Stanzl SE, Plasser O, Tschegg EK, Vasude´ van AK. Influence and micro-structure and load ratio on fatigue threshold behavior in 7075 aluminum alloy. Int J Fatigue 1999;21:S255– S62. [9] Fonte MA, Tschegg-Stanzl SE, Holper B, Tschegg EK, Vasudevan AK. The microstructure and environment influence on fatigue
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The effect of microstructure and environment on fatigue crack growth in 7049 aluminium alloy at negative stress ratios M. da Fonte a,∗, F. Romeiro b, M. de Freitas c, S.E. Stanzl-Tschegg d, E.K. Tschegg e, A.K. Vasude´van f a
Nautical School, 2780-572 Pac¸o de Arcos, Portugal b Instituto Polite´cnico de Leiria, ESTG, Portugal c Instituto Superior Te´cnico, Technical University of Lisbon, Lisbon, Portugal d Institute of Met and Physics, University of Agricultural Sciences, Turkenschanzstrasse 18, A-1180 Vienna, Austria e Institute of Applied and Technical Physics, Technical University of Vienna, Karlsplatz 13, A-1040 Vienna, Austria f Office of Naval Research Laboratory, Arlington, VA 22217, USA
Abstract The influence of environment and microstructure on fatigue crack growth has been investigated on a high strength 7049 aluminium alloy. This aluminium alloy was artificially aged to underaged (UA) and overaged (OA) microstructures. The heat treatment procedure was performed in order to obtain an UA and OA microstructure having the same yield strength properties, but differing in the mode of slip deformation: the UA alloy deforms by planar slip and that of the OA alloy by wavy slip. The crack growth measurements were performed in MT specimens at constant load ratios for R = 0, ⫺1, ⫺2, ⫺3 near-threshold and Paris regime in ambient air and vacuum conditions. Crack closure loads were measured in order to determine the Popen for each R ratio. Micromechanisms of near-threshold crack growth are briefly discussed for several concurrent processes involving environmentally assisted cracking with intrinsic microstructural effects. The results showed that the presence of humid air leads to a larger reduction in ⌬Kth for both the ageing conditions, but the UA specimens were superior probably because of crack branching. The role of environmental effect and microstructures near-threshold regime seems to be more significant than any mechanical contributions to the crack closure, such as plasticity, roughness, oxide, etc. 2003 Elsevier Ltd. All rights reserved. Keywords: Microstructure; 7049 and 7075 alloys; Slip mode; Near-threshold fatigue; Fatigue crack growth; Load ratio; Crack closure
1. Introduction The majority of investigations to date on the fatigue crack growth resistance of high strength aluminium alloys has focussed on crack advance in the so-called Paris regime (da / dNⱖ10⫺9 m / cycle). However, the behaviour in the near-threshold fatigue regime, which accounts for a significant portion of the crack propagation life, should deserve more attention when compared with the numerous studies concerning with crack closure effects. Designers and the industry in general need to obtain from the researcher intrinsic parameters such as the stress intensity factor range ⌬K and the
∗
Corresponding author. Fax: +351-21-4429-546. E-mail address: [email protected] (M. da Fonte).
0142-1123/$ - see front matter 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0142-1123(03)00150-6
maximum stress intensity factor Kmax to characterise the fatigue crack growth behaviour of materials under service loading conditions, microstructures and environment. They cannot be dependent on extrinsic testing parameters, such as crack closure and environmental effects, which are always very difficult to obtain. Synergetic effects of environment and loading make the understanding of the underlying mechanisms between microstructure and environment difficult. The microstructure and the environment play an important role in the fatigue crack growth resistance of Al alloys and have been investigated extensively during the last two decades [1–9]. Kirby and Beevers [10], on the other hand, clearly demonstrated that even the seemingly innocuous environment of laboratory air can lead to a marked increase in crack propagation rates, as compared to vac-
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uum, in the near-threshold fatigue regime of 7xxx series aluminium alloys. Lin and Stark [11] showed that microstructure–environment interactions at low fatigue crack growth rates may be completely different from those at higher fatigue crack growth rate levels. The concept of fatigue crack closure, introduced in the 1960s by Elber [12], has been used to explain a wide range of positive stress ratio crack propagation results (Rⱖ0), but less attention has been given to fatigue loading for negative stress ratios (R ⬍ 0). The crack closure concept has served to explain most of the related phenomena with decreasing crack growth rates, due to plastic deformation at crack wake roughness of fatigue surfaces, oxide, etc. Crack branching may likewise lead to reduced crack growth rates [13]. The measure of crack closure value has been until now a problem of difficult resolution and Vasude´ van and Sadananda [14–17] have pointed out pertinent problems concerning the subject, having reanalysed the crack closure concepts using dislocation theory. In the absence of closure, the crack tip driving force is related to the applied stress intensity, ⌬Kapp = Kmax⫺Kmin, and when closure is induced, the driving force is decreased to a smaller effective stress amplitude, ⌬Keff = Kmax⫺Kcl. They stated that for a crack propagation to occur, two driving forces, ⌬K and Kmax, are needed simultaneously, instead one. These two driving forces are intrinsic parameters of a material and they are valid for short as well as long cracks. The apparent differences in the behaviour of short or long cracks are related with residual stress. They extended these results in terms of intrinsic behaviour in contrast to the crack closure concepts. They concluded that crack closure has a small effect on fatigue crack growth behaviour, when compared with microstructural and environmental effects. Freitas and colleagues [18] have measured crack closure effects for negative stress ratios in a medium carbon steel DIN Ck 45 for a wide range of stress ratios from R = 0.7 to R = ⫺3. For a given negative R ratio, the Popen depends strongly on the maximum load Pmax. The ratio P open / P max becomes more negative for higher negative stress ratios but only for high applied maximum load levels Pmax. An example of the less importance of the crack closure phenomena when compared with the environment effect (in ambient air and vacuum) is shown in Fig. 1. If the closure due to only wedging action is assumed in the wake of a crack without corrosion influences, then the environmental effect in Ti–6V–4Al alloy data [19] should lie to the right of the vacuum curve. The air data in Fig. 1 lying to the left of vacuum indicate the important role of environment damage ahead of crack tip in the Ti–Al alloy. This is not experimentally observed when the environment results are compared to the results in a good vacuum. The arrow to the left of vacuum curve (square points) indicates the environ-
Fig. 1. Influence of environment on fatigue crack growth in Ti–6V– 4Al alloy [19].
ment’s contribution to fatigue in ambient air. The arrow to the right (dashed curve) of vacuum curve indicates the trend of the moist air curve, if the closure was significant. It has been experimentally observed that da / dN nearthreshold to higher growth rates are nearly independent of R ratios [19,20] in a high vacuum, which leads to a clear conclusion that plasticity does not induce closure as has been postulated three decades ago. Thus, one can infer that the role of environment near-threshold regime is strongly more significant than any mechanical contributions such plasticity, roughness, oxide, etc. Keeping in mind these features, one may conclude that ⌬Kth decreasing with R is an intrinsic fatigue property of the material for that environment. This work presents a systematic fatigue crack growth data from the near-threshold region to the Paris region in a high strength Al 7049 aluminium alloy artificially aged to underaged (UA) and overaged (OA) conditions. The objective was to obtain two microstructures having approximately the same crystallographic texture and yield strength properties, but different in the slip mode. Both macroscopic and microscopic responses are compared at different R ratios and different environments, in air and vacuum at load ratios R between R = 0 and R = ⫺3. 2. Material and experimental procedure Testing material was a high strength 7049 aluminium alloy. Specimens were cut from a cold rolled sheet of a 7049 aluminium alloy with 10 mm thickness. After solution heat treatment (470 °C, 30 min) of the cold rolled plates, one part of the material was underaged at 117 °C for 90 min and the other part was overaged for 8 h at 107 °C plus for 65 h at 163 °C. The chemical composition is
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Table 1 Composition (wt.%) of Al 7049 alloy Zn
Mg Cu
Cr
7.1
2.8
0.06 0.3
1.7
Fe
Si
Mn
Ti
Ga
0.1
0.06
0.05 0.01
Zr
Al
0.1
Bal.
listed in Table 1. The UA and OA heat treatments were designed to produce approximately the same yield strength, in order to eliminate differences in fatigue crack growth behaviour due to the possible strength effects. The yield strengths of both the microstructures are ⬇440 MPa and the ultimate tensile strength of the UA material is 16% higher and its ductility is 100% higher than that of the OA alloy (Table 2). The grain size is up to several hundred micrometres in longitudinal and a few micrometres in transverse direction. Transmission electron micrographs of the different ageing conditions are shown in Fig. 2 [4] for AL 7075 aluminium alloy closed to this Al 7049. The UA material exhibits extremely fine grain precipitate (GP) zones and h⬘ intermetallic precipitates, whereas the OA structure contains predominantly coarse h as well as h⬘ precipitates. h precipitates may be seen at grain boundaries. Tests were carried out in two different servo-hydraulic testing machines, one for testing in vacuum (Schenck, 20 kN) and another one for testing in ambient air (INSTRON, 100 kN). The crack growth measurements were performed in MT specimens (150 × 40 × 5 mm) at constant load ratios ranging from R = 0 to R = ⫺3. The test frequency was 15 Hz. Crack growth rates and ⌬K were obtained according to the standard test method for measurement of fatigue crack growth rates (ASTM E 647—1999). The central notch was 10 mm long and was made by electrical-discharge machining with h = 0.25 mm. A fatigue precracking was conducted till the crack length was approximately 1 mm for each side of the notch. The thickness of the specimen was chosen due to practical reasons related to the machine load capacity of servo-hydraulic Schenck machine (20 kN). However, we did not find any buckling problems under compressive load conditions. Crack growth measurements were made by either an optical microscope with a magnification of 50×, associated with a stroboscopic illumination at the frequency of testing or through an automatic measure-
Fig. 2. Microstructure of Al 7075 alloy [4]: (a) UA Al 7075 alloy, (b) OA Al 7075 alloy.
ment technique using the compliance method associated with load–displacement curves previously calibrated using the data from optical measurements. Special grips were designed and constructed in order to put the MT specimens inside the vacuum chamber, as seen in Fig. 3. The applied maximum loads for the OA and the UA aluminium alloy, testing in air and vacuum, were 5–6 and 7–8 kN, respectively.
Table 2 Mechanical properties of Al 7049 alloy Alloy/Temper
Yield strength Rp
7049-UA 7049-OA
445 441
0.2
(MPa) UTS (MPa) 578 497
1
Elongation (%)
Reduction area (%)
KIC (MPa m2)
17.2 8
19.1 23
– 32.0
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3. Results and discussion Crack growth data were calculated according to ASTM E 647—1999, which states: ⌬K ⫽ Kmax⫺Kmin for R ⬎ 0⌬K ⫽ Kmax for R
(1)
ⱕ0 Therefore: ⌬P ⫽ Pmax⫺Pmin for R ⬎ 0⌬P ⫽ Pmax for Rⱕ0 Results can be analysed as plots of da / dN as a function of ⌬K = Kmax. The calculation of ⌬K for MT specimens is obtained by the following expression: ⌬K ⫽
⌬P B
冪2Wsec 2 pa
pa
(2)
where a = 2a / W. The basic principle of crack closure concept is that a crack grows only if the crack is completely open. Therefore, the stress intensity factor range ⌬K should be reduced to an effective ⌬Keff. The Paris law can be written as Elber law [12]: da / dN ⫽ C(⌬Keff)m
(3)
This is a unified crack growth law for different stress ratios R, where ⌬Keff ⫽ Kmax⫺Kop Fig. 3. Chamber of vacuum with the MT specimen mounted.
Crack closure loads were measured with the compliance technique at test frequency, using a data acquisition system for a wide range of stress ratios 0ⱖRⱖ ⫺3. The centre notched MT specimens were machined after the heat treatments such that crack propagation was parallel to the rolling direction. The grips for the MT specimens used in the chamber of vacuum were specially designed and constructed such that the alignment was perfect and no bending could take place. The experiments were performed in ambient air at 20 °C and at approximately 50% relative humidity. The vacuum was approximately 3 × 10⫺3 Pa. Tests were carried out on a servo-hydraulic testing machine under constant amplitude load control with a 15 Hz sine wave. Threshold ⌬Kth values were defined as those stress intensities, where the maximum crack propagation rates did not exceed 10⫺11 m/cycle. Automatic data acquisition of load (kN) and centreline displacement (mm) was done at regular intervals, such that one recorded results in 2 × 2000 pairs of data points load–displacement.
(4)
To use this unified crack growth law, Kopen or Popen should be determined from experimental data for the different stress ratios, R and Kmax. The use of crack closure (open) concept allows a unique closure free curve for the crack growth rates at different stress ratios to be determined. The measured fatigue crack propagation rates (da / dN vs. Kmax) are shown in Figs. 4 and 5 for loading at
Fig. 4. da / dN vs. Kmax at R = 0, ⫺1, ⫺2, ⫺3 in air.
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Fig. 5.
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da / dN vs. Kmax at R = 0, ⫺1, ⫺2, ⫺3 in vacuum.
R = 0, ⫺1, ⫺2, ⫺3, in vacuum and in air. The calculation of ⌬K was performed according to ASTM E 647—1999: at negative stress ratios, ⌬K = Kmax. Fig. 4 shows the curves da / dN vs. Kmax for OA and UA aluminium alloy in ambient air conditions near threshold and Paris region. Fig. 5 shows the da / dN vs. Kmax for OA and UA aluminium alloy in vacuum conditions near threshold and Paris region. When comparing the two graphs, Figs. 4 and 5, we can see how significant the effect of microstructure and environment is on crack growth rates near-threshold regime. In vacuum, for both microstructures, UA and OA, thresholds are clearly higher than those in ambient air. The K thresholds and crack growth data are clearly different for OA conditions, and it is justified by the highest level of Pop for the UA alloy, as is shown in Fig. 6a–d. If we adopt for OA a compliance offset of 1%, Pmax / Pop vary from 0.25 for R = ⫺3 up to 0.5 for R = 0, while for UA, Pmax / Pop vary from 0.32 for R = ⫺3 up to 0.55 for R = 0. These figures also show that the crack opening level is similar during crack growth at stress ratios lower than R = ⫺1. It is to be noted that the opening load was determined using the compliance offset method described at ASTM E 647—1999, without filtering the data, which gives a scatter of the opening load level. Despite this scatter, the data show clearly that opening loads are higher for the stress ratio of R = 0 than for negative stress ratio for both alloys and that opening load remains approximately constant for stress ratio lower than R = ⫺1. This confirms the results of Figs. 4 and 5 when crack growth rate is constant from R = ⫺1 to R = ⫺3. It also must be remarked that at stress ratios R = ⫺2 to R = ⫺3, the ASTM practice may not be reasonable, since it states that the compliance that corresponds to the fully open crack confirmation is determined at 25% of cyclic load range on the unloading curve. In
Fig. 6. Determination of opening load (ASTM E 6∗47—1999 method) for the OA (a, c) and the UA (b, d) conditions, at R = 0, ⫺1, ⫺2, ⫺3.
these tests and using this statement, the crack may already be closed. There are no significant differences on crack growth rate for high negative R ratios (R = ⫺2, ⫺3,⫺4). The points are parallel and very close to the curves obtained on Paris region. The strong difference is for R = ⫺1 and R = 0. The measurement of P / P max vs. compliance offset shows these differences clearly for the OA. The influence of environment, i.e. humid air, is plotted for both alloys for a stress ratio of R = ⫺1, ⫺2, ⫺3 in Fig. 4. Two effects are obvious. Firstly, the thresholds are lower in the OA than in the UA conditions and the crack growth rates are higher for the same Kmax; and secondly, the slope of the Paris line is reduced in comparison to UA and the vacuum curves. Table 3 shows the effect of microstructure and environment on the slope m of the Paris law. Table 3 Slope m of Paris law at negative R ratios, in air and vacuum R ratios
da / dN = C (⌬K)m (UA) m
⫺3, ⫺2, ⫺1, 0
Air 7.2
(OA) m Vacuum 8.9
Air 5.1
Vacuum 6.2
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Fig. 7. Fracture surfaces after fatigue loading in the threshold regime. (a) OA in air, R = ⫺3; (b) OA in vacuum, R = ⫺3.
Fig. 8. Fracture surfaces after fatigue loading in the threshold regime. (a) UA in air, R = ⫺1; (b) UA in vacuum, R = ⫺1.
Figs. 7–9 show the fracture surfaces of the OA and UA 7049 aluminium alloy after loading at negative R ratios on the threshold regime in air and vacuum. The different negative R ratios did not result in different fracture surface appearance. This result corresponds to that of the identical crack growth curves. The reproduced figures therefore are considered to represent the structures of near-threshold tested specimens at all tested negative R ratios. In Fig. 7, the fracturing features of the OA alloy after loading at R = ⫺3 in air (Fig. 7a) and vacuum (Fig. 7b) are compared. Both fractures appear smooth, ductile and flat. In the UA condition, however, a more brittle fracture, with crystallographic features, is observed after loading in air (Fig. 8a). These features were eventually observed likewise in specimens, which have been loaded in vacuum, but there mostly smooth and ductile features prevail (Fig. 8b). In addition, secondary cracks, which point to crack branching, are recognised (Fig. 9a). Crack branching was indeed observed during the measurements of the UA alloy in air and is considered as an important mechanism for crack growth retardation.
Fig. 9b was taken with a tilt angle of 35° in the SEM and thus makes the high fracture surface roughness visible. The higher fracture surface roughness of the UA specimens could be another reason for the better fatigue crack growth properties of this alloy. The reported results show that the environment and the microstructure strongly influence the fatigue crack growth resistance of high strength aluminium alloys as is known from the literature. The microstructure and the applied ⌬K primarily control the slip mode for crack growth, which can be significantly changed by the presence of the environment. The apparent differences in fatigue crack growth resistance of the two ageing conditions are ascribed to a complex interaction between two mechanisms: the embrittling effect of humid air resulting in conventional corrosion fatigue process, the role of microstructure and slip mode in inducing crack deflection and crack closure arising from a combination of environment and microstructural contributions.Corrosion fatigue processes, i.e. active path corrosion and hydrogen embrittlement [11,21,22], tend to accelerate crack growth. Although such mechanisms are not clearly understood for alu-
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slope m significantly different for both microstructures (OA and UA) and different environment conditions, in air and vacuum (see Table 3).
4. Conclusions
Fig. 9. Fracture surfaces after fatigue loading in the threshold regime. (a) UA in air, R = ⫺1; (b) UA in vacuum, R = ⫺3 (tilt 35°).
minium alloys, it is generally accepted that hydrogen embrittlement induced by water vapour with freshly formed fracture surface is responsible for the resulting enhancement of crack propagation rates in comparison with an active path dissolution mechanism. The results of this study show that the presence of humid air leads to a significant reduction of Kmax, Kth for both ageing conditions (OA, UA) when compared with the testing vacuum conditions (3 × 10⫺3 Pa). There is no evidence for the oxide induced closure. The apparent differences in fatigue crack growth resistance of the two microstructures are ascribed to a complex interaction between several mechanisms such as change of fracture mode due to different microstructure and slip mode, as well as hydrogen embrittlement. The different microstructures manifested in terms of their different deformation slip mode of planar vs. wavy indicate that the resistance to crack growth in a planar slip alloy is significantly improved in comparison to that of the wavy slip alloy, with crack branching playing an important role in the underaged alloy. Another result concerns the C and m parameters of the Paris law, which are traditionally considered as characteristic material properties. The results showed a
The fatigue crack growth threshold behaviour of an Al 7049 alloy was studied by using a material with identical chemical composition and yield strength, but different microstructure. Experiments at negative R ratios and different environment (ambient air and vacuum) showed that Kmax thresholds depend on the different microstructure and the associated deformation mechanisms. These are homogeneous slip in the OA condition and localised slip with crystallographic cracking and with a tendency to crack branching in the UA material. The second most important influence comes from the environment. It shows that the Kmax threshold is reduced by approximately 50%, probably by hydrogen embrittlement. The underaged and overaged microstructures of this alloy show significant differences in the fatigue crack growth resistance at near-threshold crack growth rates. There are no significant differences on crack growth rates for higher negative R ratios either in ambient air or in vacuum. The significant difference appears between R = 0 and R = ⫺1, which is demonstrated by the crack opening load shown in the P / Pmax compliance offset. The parameters C and m of the Paris law, traditionally considered as a specific property of materials, significantly differ for each microstructure and environment.
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