- Email: [email protected]
Elevated temperature short crack fatigue behaviour in near eutectic Al–Si alloys
International Journal of Fatigue 25 (2003) 863–869 www.elsevier.com/locate/ijfatigue
Elevated temperature short crack fatigue behaviour in near eutectic Al–Si alloys M.R. Joyce a,∗, C.M. Styles b, P.A.S. Reed b a
University of Southampton, School of Engineering Sciences, Materials Research Group, Highfield, Southampton SO17 1BJ, UK b MSS, QinetiQ, Farnborough, UK
Abstract This paper considers two candidate automotive piston alloys and highlights the influence of microstructural features on fatigue behaviour. Fatigue initiation and subsequent short crack growth was assessed at 20, 200 and 350 °C. It is shown that both temperature and test frequency have a strong influence on the fatigue performance of the materials tested. The microstructure was quantitatively characterised in terms of the primary Si distribution. Together with post failure analysis, this allowed identification of critical microstructural features affecting both fatigue crack initiation and early growth. Large primary Si particles were found to act as preferential initiation sites by cracking or decohesion (dependent on test temperature) and are also sought out preferentially during short crack growth. 2003 Elsevier Ltd. All rights reserved. Keywords: Fatigue; Elevated temperature; Microstructure; Aluminium–silicon
1. Introduction In order to meet increasingly demanding legislation, the efficiency of modern small automotive engines is constantly rising. This is partly achieved through weight reduction, which coupled with design changes to reduce exhaust emissions has placed greater demands on the material performance of many key engine components. Near eutectic Al–Si alloys are used extensively as piston materials in modern, small automotive engines. This class of alloy exhibits complex multiphase microstructures, comprising primary (blocky) and eutectic (acicular) Si, Al dendrites and numerous intermetallic particles. Automotive pistons operate under complex mechanical stresses and over a wide temperature range. Through the addition of alloying elements, it is hoped to optimise high temperature fatigue performance without compromising room temperature behaviour. The room temperature fatigue behaviour for these classes of alloy has been studied extensively and has identified porosity as a critical feature [1–3]. However improved casting
Corresponding author. Tel.: +1-44-2380-592443; fax: +1-442380-593016. E-mail address: [email protected] (M.R. Joyce). ∗
0142-1123/$ - see front matter 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0142-1123(03)00157-9
techniques have reduced porosity considerably, therefore the influence of Si [4–7] and intermetallic particles [8] on fatigue behaviour has become more critical. To optimise materials design it is necessary to fully understand the fundamental fatigue behaviour of these materials at the elevated temperatures seen in service. The fatigue behaviour has been assessed at 20 °C (baseline behaviour) 200 °C (typical gudgeon pin boss temperature) and 350 °C (typical combustion chamber bowl edge temperature). Two materials are considered; an existing Federal-Mogul Al–Si alloy designated alloy 1 and a new alloy designated alloy 2.
2. Materials The materials used in the study were high performance casting alloys (Al–Si–Cu–Ni–Mg). Both materials have similar Si content (near eutectic), however alloy 2 has higher additions of Cu, Ni and Mg compared to alloy 1. Samples were extracted directly from the crown of as cast pistons to ensure representative microstructures were tested. All samples (whether for microstructural investigation or fatigue testing) were soaked at 260 °C for 100 h before analysis to provide a practical simulation of the effect of long-term high temperature service
864
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
in an engine. Fig. 1 shows the microstructure of both alloys after etching in 0.5% HF for 25 s. Both materials were seen to exhibit complex multiphase microstructures with both primary ‘blocky’ Si and eutectic acicular Si phases present. In addition numerous intermetallics were observed, identified through reference to previous work by Daykin [9]; alloy 1 was found to contain: Al9NiFe, Mg2Si, Al8FeMg3Si6 and AlFeMnSi, whilst alloy 2 contained: Al3(Cu,Ni)2, Al9NiFe, Mg2Si, and Al8FeMg3Si6. 3. Experimental Tensile tests were performed on ‘dog-bone’-type specimens at a strain rate of 0.003 min⫺1 at 20, 200 and 350 °C. Values of Young’s modulus (E), 0.2% proof stress, ultimate tensile strength (UTS) and elongation to failure (%El) were obtained from the stress/strain plots. Short crack fatigue tests were carried out using an Instron 8872 servo-hydraulic equipped with a Genlab furnace. Nominal stresses were calculated from applied loads and sample geometries using simple elasto-plastic finite element modelling. The fatigue tests were carried out in three batches as follows: 앫 Tests with periodic interruptions for crack monitoring were carried out on both alloys at 20 °C using 12 × 12 × 70 mm plain bend bars in a three-point bend loading geometry, shown in Fig. 2. Testing was carried out with a maximum top surface stress of 130% of each alloy’s 0.2% proof stress, at a load ratio of 0.1 and a frequency of 10 Hz. Prior to testing the sample top surface was highly polished to allow crack monitoring via acetate replication. Fatigue crack growth curves were constructed using Scott and Thorpe’s [10] analysis to calculate ⌬K levels. 앫 S-N type testing, without interruptions for crack monitoring, was carried out under ambient conditions using sub size (6×6×50 mm) samples due to the limited supply of material. These tests were carried out using the same loading geometry, load ratio and frequency as previously. Tests were carried out over a
Fig. 1.
Fig. 2. Schematic of room temperature short crack fatigue sample and loading geometry.
range of load levels corresponding to nominal top surface stresses of 120–170% of each alloy’s 0.2% proof stress. 앫 Preliminary samples tested at elevated temperature using the above plain bend bar format failed via gross sample deformation rather than fatigue. This strain accumulation could have been overcome by employing fully reversed loading, but unfortunately this was not possible with the available testing equipment. Instead subsequent samples were notched and the testing frequency increased to 50 Hz, thereby increasing the local stresses and reducing the time for stress relaxation to occur. The notch was 4 mm deep, with a root radius of 2 mm as shown in Fig. 3. Prior to testing the notch root was polished to a 0.25 µm finish using dental felts, allowing monitoring of fatigue crack initiation and subsequent early growth. It was necessary to periodically cool samples to room temperature to allow surface replication. This was achieved by air quenching, after replication the sample was reheated and allowed to soak for 20 min before the test was restarted. Post test analysis included optical and scanning elec-
Optical micrographs showing the general microstructure of (a) alloy 1 and (b) alloy 2 after etching in 0.5% HF for 25 s.
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
Fig. 3. Schematic of elevated temperature notched fatigue specimen and three-point bend loading geometry.
tron microscopy using a JEOL 6500F field emission gun scanning election microscope (FEGSEM) to identify critical microstructural features evident along fatigue crack paths.
4. Results 4.1. Tensile testing Results of the tensile testing are shown in Table 1. As testing temperature was increased proof stress and UTS decreased sharply and elongation increased. Across the testing temperature range alloy 2 showed slightly higher strengths but reduced ductility compared to alloy 1. 4.2. Ambient temperature short crack fatigue test results At room temperature, the first small fatigue crack was observed in alloy 1 after only 5000 cycles (N / N f = 0.75%), the overall fatigue lifetime was ~660 000 cycles. This crack was seen to have initiated adjacent to a large irregularly shaped cracked Si particle, which is shown
865
Fig. 4. Optical micrograph of fatigue crack initiation from cracked primary Si in alloy 1 at room temperature.
in Fig. 4. Twelve secondary cracks were also observed, the initiation points of which were all associated with cracked Si particles. In alloy 2, the first small fatigue crack was observed after only 1000 cycles (N / N f = 0.83%), whilst the overall fatigue life was 120 000 cycles. Again the initiation point was associated with a cracked large primary Si particle (~50 µm diameter). Of the 11 secondary cracks, nine initiated from cracked primary Si, whilst the remaining two appeared to be associated with clusters of Al3(Cu,Ni)2 intermetallics. Examples of these two microstructural features are shown in Figs. 5 and 6 respectively. The cracking of a primary Si particle does not always indicate the impending initiation of a fatigue crack. Indeed, in both alloys, the majority of cracked particles did not initiate cracks. No partially cracked primary Si particles were observed and it appears that Si cracking occurs either instantaneously or within an extremely limited number of cycles. The periodic interruptions for crack monitoring did not appear to effect fatigue behaviour at ambient temperature, lifetimes of interrupted and uninterrupted tests being similar in both alloys.
Table 1 Tensile properties Material Temperature Young’s (°C) modulus (GPa)
0.2% proof UTS stress (MPa) (MPa)
%El. (%)
Alloy Alloy Alloy Alloy Alloy Alloy
124 98 45 139 98 55
1.5 3.5 10.5 0.9 2.7 7.0
1 1 1 2 2 2
RT 200 350 RT 200 350
72 – – 73 – –
198 140 78 202 148 79
Fig. 5. Optical micrograph of fatigue crack initiation from cracked primary Si in alloy 2 at room temperature.
866
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
Fig. 6. Optical micrograph of fatigue crack initiation at a cluster of Al3(Cu,Ni)2 particles in alloy 2 at room temperature.
5. Elevated temperature short crack fatigue test results
Fig. 8. Secondary electron micrograph of fatigue crack initiation from decohered Si in alloy 2 at 350 °C.
5.1. S-N type fatigue test results At elevated temperature considerable surface oxide build up was noted, this unfortunately obscured much of the microstructure in the later portion of the fatigue lifetime. In contrast to the results at ambient temperature, fatigue lifetimes of tests interrupted for crack monitoring were considerably shorter than equivalent uninterrupted tests. Significant numbers of cracked Si particles were observed after the first thermal cycle (whether this was at 25 000 or 125 000 cycles) and the subsequent lifetime to failure after the first thermal cycle was within ~40 000–60 000 cycles. In common with the room temperature observations, fatigue crack initiation at elevated temperature was seen to be associated with the primary Si phase. However, rather than being exclusively associated with cracked Si, fatigue cracks were also seen to initiate from decohesion of the Si–matrix interface. Examples of both forms of initiation site are shown in Figs. 7 and 8 respectively.
Fig. 7. Secondary electron micrograph of fatigue crack initiation from fractured Si in alloy 2 at 350 °C (apparent decohesion thought to be a secondary effect).
Instead of plotting S-N data, Fig. 9 shows fatigue lifetimes of all the uninterrupted tests plotted against notch root plastic strain range, to better account for the localised plastic flow. The required strain values were calculated using an elasto-plastic finite element model. However, it should be noted that, in the absence of any stabilised cyclic stress–strain properties, the strain ranges are calculated using monotonic material properties. As such, it is likely that the elasto-plastic strain ranges predicted at the higher testing temperatures are overestimates, since no account is made of cyclic hardening or ratchetting effects. Considering the fatigue lifetimes of the uninterrupted tests, alloy 1 showed apparently greater fatigue resistance at all testing temperatures. 5.2. Summary of fatigue crack growth rates The room temperature da/dN vs. short crack ⌬K curves for both alloys are shown in Fig. 10 together with
Fig. 9.
Fatigue lifetimes of all uninterrupted fatigue tests.
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
867
6. Discussion 6.1. Fatigue crack initiation
Fig. 10. Comparison of long and short crack fatigue crack growth rates in both alloys at room temperature.
appropriate long crack results [11]. It can be seen that both alloys exhibit classical short fatigue crack behaviour with the crack often retarding severely, such that the crack growth between periodic interruptions was below the resolution of the monitoring method (~1 µm). These periods of reduced crack growth rate were observed when the crack tip impinged on primary Si secondary phases. Furthermore, crack growth is observed at ⌬K levels below the projected long crack thresholds, therefore fatigue life predictions calculated using the long crack data would be an overestimate if growth of short cracks dominated service lifetime. Comparison of the two materials shows that at short crack lengths, alloy 2 exhibits faster crack growth rates at room temperature. Elevated temperature interrupted tests with surface replication were performed to obtain short crack growth measurements. However, as can be seen in Table 2, the replication process has a dramatic effect on fatigue lifetime.
At room temperature, of 23 fatigue cracks initiated, 21 were observed to emanate from Si particles and only two from Al3(Cu,Ni)2 intermetallics (and these only in alloy 2). At elevated temperature all initiation events were associated with large Si particles. Other studies on Al–Si casting alloys have shown fatigue cracks to initiate at persistent slip bands (PSBs) [12], porosity [1–3], Si particles [4–7] and Fe-containing intermetallics [8]. In the current study PSBs were not observed and there was no evidence of crack initiation at porosity in the one sample (alloy 1) that exhibited some angular pores up to ~100 µm diameter. It is interesting to note that fatigue cracks were generally initiated from the large, blocky Si particles rather than the acicular form. This may be a size effect (i.e. large particle is more likely to contain a flaw, re-entrant angle etc.). Only two initiation events were observed at intermetallics (alloy 2 at room temperature only) and it should be noted that these were clusters rather than single particles. Discrete Al3(Cu,Ni)2 intermetallic particles are significantly smaller than the large blocky Si particles that are observed to fracture. Bischofberger [8] found that Fe-containing intermetallics only significantly affected initiation when the Fe content was above 0.7%; in both alloys considered in this work the Fe content is below this level. Lee et al. [13] also commented that relatively large Si particles fracture much more readily than Fe-containing intermetallics and that it is indeed the Si particle morphology that controls fatigue crack growth. It is interesting to note that, whilst at room temperature, fatigue crack initiation with respect to the primary Si is exclusively associated with particle cracking. At elevated temperature Si–matrix decohesion becomes more common. This indicates that this interface may become more susceptible to void formation and
Table 2 Elevated temperature short crack fatigue test conditions Test No.
Temp (°C)
Max load (N)
Nominal max stress (MPa)
Nominal ⑀pl,max (×10⫺3)
Replication schedule (cycles)
Cycles to failure
AE160 S1 AE160 S2 AE160 S3
350 350 350
1620 1620 1620
85.5 85.5 85.5
33.0 33.0 33.0
452 110 161 096 67 031
AE160 AE160 AE413 AE413 AE413
S4 S5 S1 S2 S3
200 350 350 350 350
1620 900 1620 1620 1620
123.0 55.0 85.8 85.8 85.8
2.1 2.1 57.6 57.6 57.6
AE413 S4 AE413 S5
350 200
1620 1620
85.8 116.0
57.6 2.55
Un-interrupted 100 × 103, 125 × 103, 150 × 103 25 × 103, 50 × 103, every 1000 cycles thereafter Uninterrupted Uninterrupted Uninterrupted Uninterrupted 25 × 103, then every 5000 to 50 × 103, then every 1000 thereafter Uninterrupted Uninterrupted
979 554 ⬎16.7×106 402 145 335 696 66 984
188 906 ⬎ 5.1 × 106
868
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
thereby more likely to initiate cracks at elevated temperatures.
Table 3 Mean value and standard deviation for primary Si population samples Population sample
Mean particle area and standard deviation (µm2)
Si background Si associated with fatigue crack initiation Si evident on short crack path
324 ± 210 1124± 574
6.2. Effect of elevated temperature At elevated temperatures, several effects were noted. Firstly it was found that a stress concentrating feature, higher loads and an increased testing frequency compared to room temperature, were all required to cause failure through fatigue rather than via gross sample deformation. This is due to an effective stress relaxation mechanism operating at elevated temperature reducing the magnitude of the stress field in the notch root. It is known that this class of material creeps readily at elevated temperatures [14] and hence to truly quantify the stress state in the notch root would require finite element modelling employing stabilised stress–strain data. In the absence of such data the elevated-temperature, plastic strain values in Fig. 9 (calculated using monotonic properties) must be viewed as overestimates. Secondly, evidence of severe sensitivity to thermo-mechanical fatigue was seen in the considerable reduction in fatigue lifetime brought about by the thermal cycling associated with the replication process. Although this precluded meaningful short crack growth measurements being obtained to compare with the room temperature case, it is noteworthy that the pistons manufactured from these materials are routinely subjected to thermal cycling in service. The large numbers of cracked Si particles observed after the first thermal cycle indicates that the thermal cycle promoted Si cracking and thereby fatigue crack initiation. The thermal expansion coefficient difference between Si and Al indicates that thermal cycling is likely to lead to cracking/decohesion of larger Si particles due to significant strain mismatch, indicating a potent crack initiation mechanism under thermo-mechanical fatigue conditions for these alloys. 6.3. Statistical study of fatigue crack initiation sites Sections of microstructure from untested samples of both alloys were randomly selected and the primary Si particles measured to provide a background population sample (~600 particles) against which to compare the population sample of initiating Si particles. Due to the limited number of identifiable initiating Si particles (34 total) no attempt was made to partition the initiating population sample according to material or testing temperature. Mean values and standard deviation of particle size for the two population samples are given in Table 3, whilst frequency histograms are shown in Fig. 11. It is clear that the particles associated with fatigue crack initiation are on average far larger than the background population sample, in terms of both mean value and a
673 ± 327
Fig. 11. Frequency distribution of Si particle size in background, initiating and path impinging populations.
marked shift in the overall population sample distribution. 6.4. Propagation behaviour The short crack growth rates shown in Fig. 10 show classical behaviour with periods of crack acceleration and arrest. In common with the work of Shiozawa et al. [2], it was shown that the periods of arrest corresponded to the crack tip impinging on primary Si particles. To investigate crack path preferentiality with respect to the primary Si population sample, a line counting approach was adopted. Micrographs of both materials were overlaid with a series of randomly orientated lines. This provided an average number of particle impingements along a line of given length. This figure was then compared with the average number of Si particles along short fatigue cracks in both materials. It was found that in both materials Si impingement was ~30% higher along short fatigue crack paths than along a random line. Furthermore, it was also noted that the mean size of primary Si on the crack path was generally larger than average (shown in Table 3 and Fig. 11), this is similar to the findings of Dighe and Gokhale [15] in cast Al–Si–Mg alloys. It therefore appears that large primary Si particles are important in controlling short crack propagation behaviour, as well as initiation.
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
7. Conclusions 앫 Alloy 1 appears to be generally more fatigue-resistant than alloy 2. 앫 Fatigue crack initiation in both alloys was shown to be associated with the primary Si phase, with only two occurrences of initiation at intermetallic clusters in alloy 2 at room temperature. Ambient temperature initiation from Si was exclusively caused by particle fracture, whereas at high temperature Si/matrix interface decohesion was also observed. In both cases, primary Si particles associated with fatigue initiation are generally larger than the population sample average. 앫 At room temperature classical short crack growth behaviour was observed, periods of decreased crack growth rate were linked to the crack tip impinging on secondary phases. A line counting approach was employed to show that primary Si is preferentially picked out by the propagating short cracks. It was also shown that these Si particles were generally larger than the background population sample. 앫 Crack growth measurements at elevated temperature were not possible since the thermal cycling associated with the replication process dramatically affected the fatigue behaviour in both alloys. 앫 In order to optimise this microstructure for fatigue performance, it appears necessary to reduce the presence of large primary Si, since this appears to be deleterious to both fatigue crack initiation and to subsequent propagation. However in doing this, intermetallic distribution may become more important in determining fatigue performance.
869
Acknowledgements The authors would like to thank EPSRC (grant GR/M38667) and Federal Mogul for financial and material support throughout the duration of this project. References [1] Verdu C, Cercueil H, Communal S, Saintfort P, Fougeres R. Materials Science Forum 1996;217-222:1449. [2] Shiozawa K, Tohda Y, Sun S-M. Fatigue and Fracture of Engineering Materials and Structures 1997;20:237. [3] Skallerud B, Iveland T, Harkegard G. Engineering Fracture Mechanics 1993;44:857. [4] Hoskin GA, Provan JW, Gruzleski JE. Theoretical and Applied Fracture Mechanics 1988;10:27. [5] Madelaine-Dupuich O, Stolarz J. Materials Science Forum 1996;217-222:1343. [6] Plumtree A, Schafer S. In: The behaviour of short fatigue cracks. London: Mechanical Engineering Publications; 1986. p. 215. [7] Inguanti PC. In: Proceedings of the 17th National SAMPE Technical Conference. 1985. p. 61. [8] Bischofberger U, Neite G, Exner HE. Key Engineering Materials 1990;44-45:333. [9] Daykin CRS. Microstructural modelling of commercial Al–Si alloys for piston applications. PhD thesis, Cambridge University, 1998. [10] Scott PM, Thorpe TW. Fatigue and Fracture of Engineering Materials and Structures 1981;4:291. [11] Styles CM, Reed PAS. Fatigue of an Al–Si gravity die casting. In: Starke EA, Sanders TH, Cassada WA, editors. Proc ICAA-7, Mater Sci Forum, vol. 331–337. Aluminium Alloys: Their Physical and Mechanical Properties. 2000. p. 1457–62. [12] Odegaard JA, Hafas JE, Pedersen K. In: Proceedings of Fatigue ’90, vol. 1. 1990. p. 273. [13] Lee FT, Major JF, Samuel FH. Metall Mater Trans A 1995;26:1553. [14] Barnes S, Shultz J. Private communication, Federal Mogul Technology, Cawston, UK. [15] Dighe MD, Gokhale AM. Scripta Materialia 1997;37:1435.
Elevated temperature short crack fatigue behaviour in near eutectic Al–Si alloys M.R. Joyce a,∗, C.M. Styles b, P.A.S. Reed b a
University of Southampton, School of Engineering Sciences, Materials Research Group, Highfield, Southampton SO17 1BJ, UK b MSS, QinetiQ, Farnborough, UK
Abstract This paper considers two candidate automotive piston alloys and highlights the influence of microstructural features on fatigue behaviour. Fatigue initiation and subsequent short crack growth was assessed at 20, 200 and 350 °C. It is shown that both temperature and test frequency have a strong influence on the fatigue performance of the materials tested. The microstructure was quantitatively characterised in terms of the primary Si distribution. Together with post failure analysis, this allowed identification of critical microstructural features affecting both fatigue crack initiation and early growth. Large primary Si particles were found to act as preferential initiation sites by cracking or decohesion (dependent on test temperature) and are also sought out preferentially during short crack growth. 2003 Elsevier Ltd. All rights reserved. Keywords: Fatigue; Elevated temperature; Microstructure; Aluminium–silicon
1. Introduction In order to meet increasingly demanding legislation, the efficiency of modern small automotive engines is constantly rising. This is partly achieved through weight reduction, which coupled with design changes to reduce exhaust emissions has placed greater demands on the material performance of many key engine components. Near eutectic Al–Si alloys are used extensively as piston materials in modern, small automotive engines. This class of alloy exhibits complex multiphase microstructures, comprising primary (blocky) and eutectic (acicular) Si, Al dendrites and numerous intermetallic particles. Automotive pistons operate under complex mechanical stresses and over a wide temperature range. Through the addition of alloying elements, it is hoped to optimise high temperature fatigue performance without compromising room temperature behaviour. The room temperature fatigue behaviour for these classes of alloy has been studied extensively and has identified porosity as a critical feature [1–3]. However improved casting
Corresponding author. Tel.: +1-44-2380-592443; fax: +1-442380-593016. E-mail address: [email protected] (M.R. Joyce). ∗
0142-1123/$ - see front matter 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0142-1123(03)00157-9
techniques have reduced porosity considerably, therefore the influence of Si [4–7] and intermetallic particles [8] on fatigue behaviour has become more critical. To optimise materials design it is necessary to fully understand the fundamental fatigue behaviour of these materials at the elevated temperatures seen in service. The fatigue behaviour has been assessed at 20 °C (baseline behaviour) 200 °C (typical gudgeon pin boss temperature) and 350 °C (typical combustion chamber bowl edge temperature). Two materials are considered; an existing Federal-Mogul Al–Si alloy designated alloy 1 and a new alloy designated alloy 2.
2. Materials The materials used in the study were high performance casting alloys (Al–Si–Cu–Ni–Mg). Both materials have similar Si content (near eutectic), however alloy 2 has higher additions of Cu, Ni and Mg compared to alloy 1. Samples were extracted directly from the crown of as cast pistons to ensure representative microstructures were tested. All samples (whether for microstructural investigation or fatigue testing) were soaked at 260 °C for 100 h before analysis to provide a practical simulation of the effect of long-term high temperature service
864
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
in an engine. Fig. 1 shows the microstructure of both alloys after etching in 0.5% HF for 25 s. Both materials were seen to exhibit complex multiphase microstructures with both primary ‘blocky’ Si and eutectic acicular Si phases present. In addition numerous intermetallics were observed, identified through reference to previous work by Daykin [9]; alloy 1 was found to contain: Al9NiFe, Mg2Si, Al8FeMg3Si6 and AlFeMnSi, whilst alloy 2 contained: Al3(Cu,Ni)2, Al9NiFe, Mg2Si, and Al8FeMg3Si6. 3. Experimental Tensile tests were performed on ‘dog-bone’-type specimens at a strain rate of 0.003 min⫺1 at 20, 200 and 350 °C. Values of Young’s modulus (E), 0.2% proof stress, ultimate tensile strength (UTS) and elongation to failure (%El) were obtained from the stress/strain plots. Short crack fatigue tests were carried out using an Instron 8872 servo-hydraulic equipped with a Genlab furnace. Nominal stresses were calculated from applied loads and sample geometries using simple elasto-plastic finite element modelling. The fatigue tests were carried out in three batches as follows: 앫 Tests with periodic interruptions for crack monitoring were carried out on both alloys at 20 °C using 12 × 12 × 70 mm plain bend bars in a three-point bend loading geometry, shown in Fig. 2. Testing was carried out with a maximum top surface stress of 130% of each alloy’s 0.2% proof stress, at a load ratio of 0.1 and a frequency of 10 Hz. Prior to testing the sample top surface was highly polished to allow crack monitoring via acetate replication. Fatigue crack growth curves were constructed using Scott and Thorpe’s [10] analysis to calculate ⌬K levels. 앫 S-N type testing, without interruptions for crack monitoring, was carried out under ambient conditions using sub size (6×6×50 mm) samples due to the limited supply of material. These tests were carried out using the same loading geometry, load ratio and frequency as previously. Tests were carried out over a
Fig. 1.
Fig. 2. Schematic of room temperature short crack fatigue sample and loading geometry.
range of load levels corresponding to nominal top surface stresses of 120–170% of each alloy’s 0.2% proof stress. 앫 Preliminary samples tested at elevated temperature using the above plain bend bar format failed via gross sample deformation rather than fatigue. This strain accumulation could have been overcome by employing fully reversed loading, but unfortunately this was not possible with the available testing equipment. Instead subsequent samples were notched and the testing frequency increased to 50 Hz, thereby increasing the local stresses and reducing the time for stress relaxation to occur. The notch was 4 mm deep, with a root radius of 2 mm as shown in Fig. 3. Prior to testing the notch root was polished to a 0.25 µm finish using dental felts, allowing monitoring of fatigue crack initiation and subsequent early growth. It was necessary to periodically cool samples to room temperature to allow surface replication. This was achieved by air quenching, after replication the sample was reheated and allowed to soak for 20 min before the test was restarted. Post test analysis included optical and scanning elec-
Optical micrographs showing the general microstructure of (a) alloy 1 and (b) alloy 2 after etching in 0.5% HF for 25 s.
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
Fig. 3. Schematic of elevated temperature notched fatigue specimen and three-point bend loading geometry.
tron microscopy using a JEOL 6500F field emission gun scanning election microscope (FEGSEM) to identify critical microstructural features evident along fatigue crack paths.
4. Results 4.1. Tensile testing Results of the tensile testing are shown in Table 1. As testing temperature was increased proof stress and UTS decreased sharply and elongation increased. Across the testing temperature range alloy 2 showed slightly higher strengths but reduced ductility compared to alloy 1. 4.2. Ambient temperature short crack fatigue test results At room temperature, the first small fatigue crack was observed in alloy 1 after only 5000 cycles (N / N f = 0.75%), the overall fatigue lifetime was ~660 000 cycles. This crack was seen to have initiated adjacent to a large irregularly shaped cracked Si particle, which is shown
865
Fig. 4. Optical micrograph of fatigue crack initiation from cracked primary Si in alloy 1 at room temperature.
in Fig. 4. Twelve secondary cracks were also observed, the initiation points of which were all associated with cracked Si particles. In alloy 2, the first small fatigue crack was observed after only 1000 cycles (N / N f = 0.83%), whilst the overall fatigue life was 120 000 cycles. Again the initiation point was associated with a cracked large primary Si particle (~50 µm diameter). Of the 11 secondary cracks, nine initiated from cracked primary Si, whilst the remaining two appeared to be associated with clusters of Al3(Cu,Ni)2 intermetallics. Examples of these two microstructural features are shown in Figs. 5 and 6 respectively. The cracking of a primary Si particle does not always indicate the impending initiation of a fatigue crack. Indeed, in both alloys, the majority of cracked particles did not initiate cracks. No partially cracked primary Si particles were observed and it appears that Si cracking occurs either instantaneously or within an extremely limited number of cycles. The periodic interruptions for crack monitoring did not appear to effect fatigue behaviour at ambient temperature, lifetimes of interrupted and uninterrupted tests being similar in both alloys.
Table 1 Tensile properties Material Temperature Young’s (°C) modulus (GPa)
0.2% proof UTS stress (MPa) (MPa)
%El. (%)
Alloy Alloy Alloy Alloy Alloy Alloy
124 98 45 139 98 55
1.5 3.5 10.5 0.9 2.7 7.0
1 1 1 2 2 2
RT 200 350 RT 200 350
72 – – 73 – –
198 140 78 202 148 79
Fig. 5. Optical micrograph of fatigue crack initiation from cracked primary Si in alloy 2 at room temperature.
866
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
Fig. 6. Optical micrograph of fatigue crack initiation at a cluster of Al3(Cu,Ni)2 particles in alloy 2 at room temperature.
5. Elevated temperature short crack fatigue test results
Fig. 8. Secondary electron micrograph of fatigue crack initiation from decohered Si in alloy 2 at 350 °C.
5.1. S-N type fatigue test results At elevated temperature considerable surface oxide build up was noted, this unfortunately obscured much of the microstructure in the later portion of the fatigue lifetime. In contrast to the results at ambient temperature, fatigue lifetimes of tests interrupted for crack monitoring were considerably shorter than equivalent uninterrupted tests. Significant numbers of cracked Si particles were observed after the first thermal cycle (whether this was at 25 000 or 125 000 cycles) and the subsequent lifetime to failure after the first thermal cycle was within ~40 000–60 000 cycles. In common with the room temperature observations, fatigue crack initiation at elevated temperature was seen to be associated with the primary Si phase. However, rather than being exclusively associated with cracked Si, fatigue cracks were also seen to initiate from decohesion of the Si–matrix interface. Examples of both forms of initiation site are shown in Figs. 7 and 8 respectively.
Fig. 7. Secondary electron micrograph of fatigue crack initiation from fractured Si in alloy 2 at 350 °C (apparent decohesion thought to be a secondary effect).
Instead of plotting S-N data, Fig. 9 shows fatigue lifetimes of all the uninterrupted tests plotted against notch root plastic strain range, to better account for the localised plastic flow. The required strain values were calculated using an elasto-plastic finite element model. However, it should be noted that, in the absence of any stabilised cyclic stress–strain properties, the strain ranges are calculated using monotonic material properties. As such, it is likely that the elasto-plastic strain ranges predicted at the higher testing temperatures are overestimates, since no account is made of cyclic hardening or ratchetting effects. Considering the fatigue lifetimes of the uninterrupted tests, alloy 1 showed apparently greater fatigue resistance at all testing temperatures. 5.2. Summary of fatigue crack growth rates The room temperature da/dN vs. short crack ⌬K curves for both alloys are shown in Fig. 10 together with
Fig. 9.
Fatigue lifetimes of all uninterrupted fatigue tests.
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
867
6. Discussion 6.1. Fatigue crack initiation
Fig. 10. Comparison of long and short crack fatigue crack growth rates in both alloys at room temperature.
appropriate long crack results [11]. It can be seen that both alloys exhibit classical short fatigue crack behaviour with the crack often retarding severely, such that the crack growth between periodic interruptions was below the resolution of the monitoring method (~1 µm). These periods of reduced crack growth rate were observed when the crack tip impinged on primary Si secondary phases. Furthermore, crack growth is observed at ⌬K levels below the projected long crack thresholds, therefore fatigue life predictions calculated using the long crack data would be an overestimate if growth of short cracks dominated service lifetime. Comparison of the two materials shows that at short crack lengths, alloy 2 exhibits faster crack growth rates at room temperature. Elevated temperature interrupted tests with surface replication were performed to obtain short crack growth measurements. However, as can be seen in Table 2, the replication process has a dramatic effect on fatigue lifetime.
At room temperature, of 23 fatigue cracks initiated, 21 were observed to emanate from Si particles and only two from Al3(Cu,Ni)2 intermetallics (and these only in alloy 2). At elevated temperature all initiation events were associated with large Si particles. Other studies on Al–Si casting alloys have shown fatigue cracks to initiate at persistent slip bands (PSBs) [12], porosity [1–3], Si particles [4–7] and Fe-containing intermetallics [8]. In the current study PSBs were not observed and there was no evidence of crack initiation at porosity in the one sample (alloy 1) that exhibited some angular pores up to ~100 µm diameter. It is interesting to note that fatigue cracks were generally initiated from the large, blocky Si particles rather than the acicular form. This may be a size effect (i.e. large particle is more likely to contain a flaw, re-entrant angle etc.). Only two initiation events were observed at intermetallics (alloy 2 at room temperature only) and it should be noted that these were clusters rather than single particles. Discrete Al3(Cu,Ni)2 intermetallic particles are significantly smaller than the large blocky Si particles that are observed to fracture. Bischofberger [8] found that Fe-containing intermetallics only significantly affected initiation when the Fe content was above 0.7%; in both alloys considered in this work the Fe content is below this level. Lee et al. [13] also commented that relatively large Si particles fracture much more readily than Fe-containing intermetallics and that it is indeed the Si particle morphology that controls fatigue crack growth. It is interesting to note that, whilst at room temperature, fatigue crack initiation with respect to the primary Si is exclusively associated with particle cracking. At elevated temperature Si–matrix decohesion becomes more common. This indicates that this interface may become more susceptible to void formation and
Table 2 Elevated temperature short crack fatigue test conditions Test No.
Temp (°C)
Max load (N)
Nominal max stress (MPa)
Nominal ⑀pl,max (×10⫺3)
Replication schedule (cycles)
Cycles to failure
AE160 S1 AE160 S2 AE160 S3
350 350 350
1620 1620 1620
85.5 85.5 85.5
33.0 33.0 33.0
452 110 161 096 67 031
AE160 AE160 AE413 AE413 AE413
S4 S5 S1 S2 S3
200 350 350 350 350
1620 900 1620 1620 1620
123.0 55.0 85.8 85.8 85.8
2.1 2.1 57.6 57.6 57.6
AE413 S4 AE413 S5
350 200
1620 1620
85.8 116.0
57.6 2.55
Un-interrupted 100 × 103, 125 × 103, 150 × 103 25 × 103, 50 × 103, every 1000 cycles thereafter Uninterrupted Uninterrupted Uninterrupted Uninterrupted 25 × 103, then every 5000 to 50 × 103, then every 1000 thereafter Uninterrupted Uninterrupted
979 554 ⬎16.7×106 402 145 335 696 66 984
188 906 ⬎ 5.1 × 106
868
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
thereby more likely to initiate cracks at elevated temperatures.
Table 3 Mean value and standard deviation for primary Si population samples Population sample
Mean particle area and standard deviation (µm2)
Si background Si associated with fatigue crack initiation Si evident on short crack path
324 ± 210 1124± 574
6.2. Effect of elevated temperature At elevated temperatures, several effects were noted. Firstly it was found that a stress concentrating feature, higher loads and an increased testing frequency compared to room temperature, were all required to cause failure through fatigue rather than via gross sample deformation. This is due to an effective stress relaxation mechanism operating at elevated temperature reducing the magnitude of the stress field in the notch root. It is known that this class of material creeps readily at elevated temperatures [14] and hence to truly quantify the stress state in the notch root would require finite element modelling employing stabilised stress–strain data. In the absence of such data the elevated-temperature, plastic strain values in Fig. 9 (calculated using monotonic properties) must be viewed as overestimates. Secondly, evidence of severe sensitivity to thermo-mechanical fatigue was seen in the considerable reduction in fatigue lifetime brought about by the thermal cycling associated with the replication process. Although this precluded meaningful short crack growth measurements being obtained to compare with the room temperature case, it is noteworthy that the pistons manufactured from these materials are routinely subjected to thermal cycling in service. The large numbers of cracked Si particles observed after the first thermal cycle indicates that the thermal cycle promoted Si cracking and thereby fatigue crack initiation. The thermal expansion coefficient difference between Si and Al indicates that thermal cycling is likely to lead to cracking/decohesion of larger Si particles due to significant strain mismatch, indicating a potent crack initiation mechanism under thermo-mechanical fatigue conditions for these alloys. 6.3. Statistical study of fatigue crack initiation sites Sections of microstructure from untested samples of both alloys were randomly selected and the primary Si particles measured to provide a background population sample (~600 particles) against which to compare the population sample of initiating Si particles. Due to the limited number of identifiable initiating Si particles (34 total) no attempt was made to partition the initiating population sample according to material or testing temperature. Mean values and standard deviation of particle size for the two population samples are given in Table 3, whilst frequency histograms are shown in Fig. 11. It is clear that the particles associated with fatigue crack initiation are on average far larger than the background population sample, in terms of both mean value and a
673 ± 327
Fig. 11. Frequency distribution of Si particle size in background, initiating and path impinging populations.
marked shift in the overall population sample distribution. 6.4. Propagation behaviour The short crack growth rates shown in Fig. 10 show classical behaviour with periods of crack acceleration and arrest. In common with the work of Shiozawa et al. [2], it was shown that the periods of arrest corresponded to the crack tip impinging on primary Si particles. To investigate crack path preferentiality with respect to the primary Si population sample, a line counting approach was adopted. Micrographs of both materials were overlaid with a series of randomly orientated lines. This provided an average number of particle impingements along a line of given length. This figure was then compared with the average number of Si particles along short fatigue cracks in both materials. It was found that in both materials Si impingement was ~30% higher along short fatigue crack paths than along a random line. Furthermore, it was also noted that the mean size of primary Si on the crack path was generally larger than average (shown in Table 3 and Fig. 11), this is similar to the findings of Dighe and Gokhale [15] in cast Al–Si–Mg alloys. It therefore appears that large primary Si particles are important in controlling short crack propagation behaviour, as well as initiation.
M.R. Joyce et al. / International Journal of Fatigue 25 (2003) 863–869
7. Conclusions 앫 Alloy 1 appears to be generally more fatigue-resistant than alloy 2. 앫 Fatigue crack initiation in both alloys was shown to be associated with the primary Si phase, with only two occurrences of initiation at intermetallic clusters in alloy 2 at room temperature. Ambient temperature initiation from Si was exclusively caused by particle fracture, whereas at high temperature Si/matrix interface decohesion was also observed. In both cases, primary Si particles associated with fatigue initiation are generally larger than the population sample average. 앫 At room temperature classical short crack growth behaviour was observed, periods of decreased crack growth rate were linked to the crack tip impinging on secondary phases. A line counting approach was employed to show that primary Si is preferentially picked out by the propagating short cracks. It was also shown that these Si particles were generally larger than the background population sample. 앫 Crack growth measurements at elevated temperature were not possible since the thermal cycling associated with the replication process dramatically affected the fatigue behaviour in both alloys. 앫 In order to optimise this microstructure for fatigue performance, it appears necessary to reduce the presence of large primary Si, since this appears to be deleterious to both fatigue crack initiation and to subsequent propagation. However in doing this, intermetallic distribution may become more important in determining fatigue performance.
869
Acknowledgements The authors would like to thank EPSRC (grant GR/M38667) and Federal Mogul for financial and material support throughout the duration of this project. References [1] Verdu C, Cercueil H, Communal S, Saintfort P, Fougeres R. Materials Science Forum 1996;217-222:1449. [2] Shiozawa K, Tohda Y, Sun S-M. Fatigue and Fracture of Engineering Materials and Structures 1997;20:237. [3] Skallerud B, Iveland T, Harkegard G. Engineering Fracture Mechanics 1993;44:857. [4] Hoskin GA, Provan JW, Gruzleski JE. Theoretical and Applied Fracture Mechanics 1988;10:27. [5] Madelaine-Dupuich O, Stolarz J. Materials Science Forum 1996;217-222:1343. [6] Plumtree A, Schafer S. In: The behaviour of short fatigue cracks. London: Mechanical Engineering Publications; 1986. p. 215. [7] Inguanti PC. In: Proceedings of the 17th National SAMPE Technical Conference. 1985. p. 61. [8] Bischofberger U, Neite G, Exner HE. Key Engineering Materials 1990;44-45:333. [9] Daykin CRS. Microstructural modelling of commercial Al–Si alloys for piston applications. PhD thesis, Cambridge University, 1998. [10] Scott PM, Thorpe TW. Fatigue and Fracture of Engineering Materials and Structures 1981;4:291. [11] Styles CM, Reed PAS. Fatigue of an Al–Si gravity die casting. In: Starke EA, Sanders TH, Cassada WA, editors. Proc ICAA-7, Mater Sci Forum, vol. 331–337. Aluminium Alloys: Their Physical and Mechanical Properties. 2000. p. 1457–62. [12] Odegaard JA, Hafas JE, Pedersen K. In: Proceedings of Fatigue ’90, vol. 1. 1990. p. 273. [13] Lee FT, Major JF, Samuel FH. Metall Mater Trans A 1995;26:1553. [14] Barnes S, Shultz J. Private communication, Federal Mogul Technology, Cawston, UK. [15] Dighe MD, Gokhale AM. Scripta Materialia 1997;37:1435.