Fatigue Life Assessment of A356 Aluminum Alloy used for Engine Cylinder Head

doi:10.21203/rs.3.rs-2756317/v1

The main aim of this study is assessing fatigue resistance of casting A356 alloy for the manufacture of engine cylinder heads at high temperatures. Testing for isothermal and in-phase thermo-mechanical fatigue was conducted at 120ºC and 280ºC and from 120ºC to 280ºC, respectively. The fatigue crack growth rate was estimated, employing different waveforms under displacement and load control. The objective to perform the relaxation tests at temperatures ranging from 75^oC to 280^oC was to obtain further information for modeling. The results showed that casting defects and heterogeneities of the material in conjunction with long testing times at high temperatures and under different loading waveforms are critical factors in determining fatigue performance. These research findings will provide parameters for simulating the evolution of mechanical properties of the material under investigation under cyclic loadings at high temperatures, whether isothermal or thermo-mechanical, to later validate the crack propagation model proposed by Pommier and Risbet (2005).

cast-aluminum

thermo-mechanical fatigue

fatigue crack growth

high temperature

engine block

Engine parts usually are subjected to temperature or/and loading variations during service [1]. Parts subjected to high temperatures display numerous problems, such as creep, oxidation (or corrosion), and fatigue. These damage mechanisms can act independently or in conjunction (which complicates the study of damage at high temperatures), according to variables and service conditions of a given material, e.g., maximum and minimum temperatures, temperature range, mechanical strain or stress ranges, strain rate (in-phase or out-of-phase) or environmental factors [2–4]. High temperatures and mechanical loadings are conducive to crack nucleation and propagation and, as a result, fatigue failure, which may be of a high or low cycle and occur isothermally or anisothermally [5]. Therefore, knowledge of both low-cycle fatigue resistance and factors controlling creep-fatigue interaction is vital to the development of components capable of operating at limit conditions.

Sehitoglu (1996) and Sarabanda (1991) showed a component under thermal cycling with mechanical loading, it will undergo cyclic thermal creep (thermal ratcheting), i.e., a gradual accumulation of inelastic deformation at the end of each cycle [1,6]. This type of failure involves ductile rupture, which sometimes unintentionally occurs in a component during service when parts of it are subjected to heat and become warmer than its surroundings, thereby causing lumps to appear.

Another aggravating factor is the presence of porosities, which have a catastrophic effect on mechanical properties of cast aluminum alloys. Just 1% volumetric fraction of pores can reduce fatigue life by 50% and reduce the strength limit by 20% as compared to the same alloy without pores [7]. For aluminum alloys, fatigue life increases as stress amplitude decrease, and the number of cycles to failure remarkably decreases as porosity fraction increases for any given stress level [8]. Fatigue life uncertainty increases as stress and/or amount of porosity decreases, which indicates the importance of pore size and fatigue life distribution.

When a combustion engine starts, the temperature of its cylinder head and valves increases rapidly at the same time that the surrounding material, which is water-cooled, keeps the thermal expansion of the material in check. Therefore, when heated, the component surface is compressed by the fresher material below its surface, which is subject to compressive stresses. Under cooling, strain occurs in the opposite direction, and tensile stresses may develop. The damage on the surface by Thermal Fatigue occurs during heating and cooling cycles, which generate temperature gradients and, as a result, internal stresses [1]. Also, mechanical strains derived from external loads generated by fuel explosion or by constraints complicate matters even further. In this case, there occurs a phenomenon is known as Thermo-mechanical Fatigue, i.e., fatigue under variation of both temperature and mechanical strain. This fatigue may give rise to several damage mechanisms that affect the deformation characteristics of the material, which are inexorably associated with damage processes [9]. Hence, due to repeated engine startup-shutdown, the whole phenomena can be treated as thermo-mechanical fatigue.

This study aimed to investigate the fatigue behavior of A356 Al alloy, under isothermal and in-phase thermo-mechanical fatigue conditions, fatigue crack propagation rate as function of temperature and waveform, and relaxation tests for future application to a computational model proposed by Pommier and Risbet (2005) to predict the fatigue crack growth rate at high temperature [10–12].

According to ASTM E3-11, the samples receiving the usual metallography preparation etched with Keller’s reagent, to assess the microstructure of α-phase size and morphology, eutectic phase distribution, and porosity using optical and scanning electron microscopes coupled to quantitative image analysis software [16, 17]. Figure 1a shows a specimen from the cylinder head to performs the mechanical tests; Figure 1b shows the specimens for tensile, isothermal fatigue, in-phase thermo-mechanical fatigue, relaxation, and fatigue crack propagation tests.

Figure 2 presents specimen size and geometry for tensile, relaxation and fatigue tests.

Tensile tests were conducted at room temperature, 120°C, and 280°C according to ASTM E21-09 and ASTM E8/E8M-13, under displacement control at 0.003 mm/s [18, 19].

The testing for isothermal fatigue under controlled strain following ASTM E 606-12, 1Hz frequency, mechanical deformation ratio R_e = -1, and at two fixed temperatures: 120°C and 280°C (the starting and final temperatures of in-phase thermo-mechanical fatigue) [20]. The reason for the temperatures is because these values usually reached within engine cylinder heads. The failure criterion was a decreasing of 50% of the maximum stress which reached in the tests. Testing for in-phase thermo-mechanical fatigue was carried out according to ASTM E2368-10, from 120°C to 280°C, similar R_e, and in relatively long periods (about 300s), i.e., a lower frequency (0.003Hz), which is quite different from that used for estimating isothermal fatigue because is difficulty in cooling down specimens while avoiding excessive temperature gradients [21]. Besides, testing for in-phase thermo-mechanical fatigue was conducted at a temperature range in a triangular waveform, with the highest temperature coinciding with maximum mechanical deformation and vice-versa with both minimum values.

Corner notch-type specimens, shaped as per Pickard (1986), were tested for crack propagation at fixed temperatures (75°C, 120°C, 200°C, and 280°C) and loadings at stress ratio R = 0.1 obtained by a sinusoidal waveform and 10 Hz frequency [22]. A trapezoidal waveform, consisting of loading and unloading ramps of 10s each and a 180s dwell at maximum loading level, producing a total of 0.005Hz frequency, was applied at the same temperatures — the monitoring of cracks occurred through EPD (electrical potential drop technique). Testing for stress relaxation was conducted in tensile at 80°C, 100°C, 120°C, 140°C, 180°C, 240°C, and 280°C for total deformation levels of 0.1%, 0.3%, 0.5%, 0.7%, and 0.9%. The tests lasted 5 minutes, under strain control.

The analysis of the morphological aspect of fracture surfaces of some specimens tested for isothermal fatigue, in-phase thermo-mechanical fatigue, and crack propagation tests through SEM (scanning electron microscopy).

Material

Table 1 shows a comparison of chemical analysis results in nominal values found in ASTM A356 [23]. Also, Table 1 indicates that Fe and Cu amounts were slightly above nominal ones.

Table 1

Chemical analysis of A356 Al alloy (wt %).
Elements	Si	Mg	Fe	Cu	Zn	Mn	Ti	Al
Nominal	6.5–8.5	0.25–0.45	0.1–0.7	0.10-3.0	0.10–0.80	0.10–0.40	0.10 _max	86.0–92.0
Actual	8.3	0.33	0.98	3.13	0.66	0.45	0.03	Base

Larger amounts of iron and copper—the most common impurities in Al-Si alloys—may be detrimental to their mechanical properties, since Cu weakens and Fe promotes the formation of acicular platelets in these alloys at high temperatures, which significantly reduce their ductility and fracture toughness. However, microstructural results in Fig. 3 indicate, as observed by Crepeau, 1998, the occurrence of a Chinese script-shaped phase, which is less detrimental to mechanical properties, due to iron forming this more compact type of eutectic phase when balanced with manganese [24].

Chemical composition analysis was performed by EDS (Energy Dispersive System) on fractured surfaces of specimens tested for fatigue crack propagation at different temperatures. It comprised the selection of 900µm² sections of the crack propagation (CP) and final fracture (FF) regions. Results from this analysis show locally embrittled regions with high Cu and Fe concentrations (Table 2).

Table 2

EDX (wt %) chemical analysis performed on the fracture surface.
Test Temperature	Region	Alloy Elements
Test Temperature	Region	Al	Si	Mg	Mn	Fe	Cu
120°C	FF*	61.27	24.08	0.12	2.05	4.41	8.07
120°C	CP**	69.72	29.23	0.33	0.03	0.01	1.4
200°C	FF*	67.41	6.80	1.66	0.18	0.29	23.67
200°C	CP **	43.68	49.43	0.33	1.62	2.26	3.34
280°C	FF*	48.47	19.02	0.14	1.61	4.91	25.85
280°C	CP**	77.37	9.88	0.93	0.10	0.13	11.59
Final Fracture; *Crack Propagation

Figure 3 shows that the microstructure comprises primary α-phase, Si-rich eutectic phase, precipitates (Mg₂Si, Al₂Cu), and inter-metallic components (Al₅FeSi, Al₈Mg₃FeSi₆, Al₅Mg₈Cu₂Si₆). The microstructure is composed of well-defined dendrite structures with secondary dendrite arm spacing (SDAS) of approximately 25µm, high porosity (Fig. 3a and 3b, dotted-dashed arrow), in addition to considerable change in silicon particles (Fig. 3b), full arrow). By ASM Handbook [17] the presence of FeMg3Si6Al8-type (Chinese-writing type) (Fig. 3(b), dashed arrow) and Fe2Si2Al9-type (blade type) precipitates (Fig. 3(b), dotted arrow).

Porosity was estimated using micrographs randomly taken in both transversal and perpendicular directions under an optical microscope coupled to an image acquisition system. The material under analysis displayed an average porosity of 9.49% in area, with typical porosity of 50µm equivalent diameter. The analysis showed some voids with up to 500µm equivalent diameter. These solidification voids act as stress concentrators and prematurely nucleate cracks inside or near the surface, thus decreasing or eliminating fatigue life for nucleation.

Mechanical Tests

Figure 4a below shows the curve of strain amplitude versus a several of reversals to fail for isothermal fatigue tests at 120°C and 280°C (0.4 of melting temperature) as well as equivalent stress amplitude versus the number of reversals to failure (Fig. 4b). It is possible to observe a sharp drop in life when test temperature was raised to 280°C, which is due to different mechanisms promoting crack nucleation as observed by Suresh (1998), e.g., cyclic slip-induced cracks, grain boundary cavitation, grain boundary sliding and wedge crack development, void nucleation and growth from both inclusions and precipitates, and oxidation and corrosion [25].

Figure 5 shows stress versus plastic deformation hysteresis curves obtained from isothermal fatigue data at 120°C and 280°C for 0.3% total mechanical deformation, which observes two distinct behaviors: A356 hardens cyclically at 120°C and softens cyclically at 280°C. As is well known, high temperatures have the effect of promoting atom rearrangement in crystal structures, i.e., atoms shift around in order to reach a better and less faulty arrangement, thus facilitating the movement of dislocations.

It is also noteworthy that under tensile and compression loadings at both 120°C and 280°C, A356 displays distinct behaviors in terms of stress magnitude. This difference in mechanical behavior is more marked at lower temperatures. At higher temperatures, this difference is subtler as dynamic micro-structural recovering is favored, which promptly decreases the density of dislocations, and the material is allowed to deform without further increase in stress plastically. Porosity, nearly 10% in this material, may have contributed to the difference observed, since it anchors dislocations.

Figure 6 shows results from in-phase thermo-mechanical fatigue tests carried out at temperatures ranging from 120°C to 280°C. Also, generates isothermal fatigue test results on the same graph for comparison. Sehitoglu (1996) observed that, in general, isothermal tests could not adequately represent thermo-mechanical fatigue behavior (TMF). Existe Aan exception when both cost and time to conduct TMF tests are taken into consideration [1]. Along these lines observes that A356 life is quite similar to that obtained for isothermal fatigue at 280°C when tested for in-phase thermo-mechanical fatigue. Besides, it may be inferred that significant damage when testing for in-phase thermo-mechanical fatigue almost exclusively occurred at maximum temperature.

Despite having tested for isothermal fatigue at a lower frequency (0.1 Hz) as compared to that employed to estimate in-phase thermo-mechanical fatigue (0.003Hz), this frequency can be considered too high. Therefore, during in-phase thermo-mechanical fatigue tests, exposition time at higher temperatures gives rise to another phenomenon, known as relaxation, which explains the difference between the hysteresis loops obtained for isothermal fatigue and in-phase thermo-mechanical fatigue, as shown by stress-strain behavior in Fig. 7. This phenomenon interacts with fatigue due to both extended testing periods and the fact that testing is under strain control. Hence, even though fatigue life is quite similar and may be used to estimate in-phase thermo-mechanical fatigue life for this Al alloy, this difference is significant, and is necessary considers the relaxation in fatigue life simulation.

To identify viscous behavior in the in-phase thermo-mechanical fatigue tests shown in Fig. 7, stress relaxation was measured at fixed temperatures of 80°C, 100°C, 120°C, 140°C, 180°C, 240°C, and 280°C, for total deformation levels of 0.1%, 0.3%, 0.5%, 0.7%, and 0.9%, and under five-minute dwell. Testing was performed under strain control, with trapezoidal waveform, using a special extensometer for measuring strain at high temperatures. Figure 8 shows results for 100°C and 280°C. Similar to the hysteresis from isothermal fatigue tests, the hysteresis obtained for relaxation tests indicates two distinct behaviors, with 240°C being the threshold temperature. A356 hardens cyclically at lower temperatures while softening cyclically at higher temperatures. High temperatures promote atom rearrangement within the crystal structure, i.e., atoms shift positions to reach a better and less faulty arrangement., which, in turn, facilitates dislocations, causing the material to soften.

Load vs. elastic and plastic deformation curves shown in Fig. 9, which also indicate that to testing under controlled strain, i.e., the total strain was kept constant throughout the test, there was an increase in plastic deformation and, consequently, a decrease in elastic deformation, so that the total deformation remained constant. Therefore, this increase in plastic deformation corresponded to creep deformation. Under 5-minute dwell time and at 0.5% total strain, plastic deformation increased and, as a result, the load necessary to keep total strain constant decreased. Therefore, given that specimens remain at high temperatures for a long time because of the meager strain rate in in-phase thermo-mechanical fatigue tests, the influence of creep is quite strong.

Notwithstanding, despite their significant size, as shown in Fig. 3, solidification voids can hardly be detected through conventional techniques, such as ultrasound and X-ray, employed regularly at the production line, especially on specimens used for microstructure analysis. More important were the observations made on the fracture surface of specimens tested for isothermal fatigue under TMF conditions; they indicated that under cyclic loading conditions, solidification voids quickly nucleated cracks in the component.

Therefore, it had based the component life on crack propagation life, and the execution of testing on specimens deriving from a cylinder head. This fact limited the size of specimens, which were machined to feature a unique geometry, as shown in Fig. 2. The temperature for the tests was 75°C, 120°C, 200°C, and 280°C, and loading ratio of R = 0.1, 10 Hz frequency, and under two different loading conditions: (a) sinusoidal waveform and (b) trapezoidal waveform (dwell), the latter consisting of loading and unloading ramps lasting 10s each and a 3 min dwell at constant load (dwell-L) and under displacement control (dwell-d) at approximately 0.005Hz. The evaluation of the crack growth according to ASTM E1457/13 by potential drop methodology was used for [26].

Figure 10 shows results from crack propagation tests. As regards fatigue tests under sinusoidal waveform and load control, for the same ΔK, as the test temperature increased, the fatigue crack growth rate gradually increased because of decreasing mechanical strength, as generally reported in the literature. The results showed a general tendency for increased spread of data as temperature rises, which can be attributed to the influence of temperature on the material microstructure, exacerbating its flaws, e.g., porosity and lack of chemical homogeneity, that have an impact on its local mechanical strength. As to the dwell type cycle, for the same temperature, it is possible to note that dwell loadings are comparatively more detrimental since they give rise to creep, which interacts with mechanical fatigue, causing crack growth rates to change significantly. Still, considering the fatigue crack propagation at 280ºC under load and displacement control, a very different fatigue behavior is found. In the case of displacement control, stress relaxation takes place at the crack front, ΔK decreases rapidly, and, as a result, the crack propagation rate is nearly constant.

Figure 11 shows displacement measurements over time during tests for fatigue crack growth rate at 120°C and 280°C, under dwell cycles and controlled load. For a constant initial load and similar initial crack size, during tests, displacement at 280°C was significantly higher, thereby indicating accumulation of creep deformation over time. This phenomenon interacts during long plateau periods, causing the material to ‘flow.’ Thus, cracks propagate due to the interaction between fatigue and creep, i.e., not only due to fatigue as in the case of sinusoidal wave loading. Creep may influence both crack initiation and growth because creep damage is caused by cavity nucleation on grain boundaries because of coalescing voids, stacking dislocations, and sliding grain boundaries. The same can happen to second-phase particles. Subsequent growth of these cavities is conducive to grain-boundary cracking and inter-granular fracture. Under static loading, failure occurs using catastrophic fracture when cavities coalesce.

Fracture surface analyses were performed using a scanning electron microscope (SEM) in conjunction with the deformation map for Al presented by Roesler et al. (2007) to guide the discussion [27]. Figure 12 presents SEM fractography from the stable crack growth region at 120ºC under sinusoidal (Fig. 12a) and dwell (Fig. 12b) cycles. In both cases, mechanical fatigue constituted the micro-mechanism, with the precise formation of striations, as shown in Fig. 12d. At 200ºC (Fig. 12c), the identification of fatigue striations on the fracture surface (Fig. 12d) since at higher temperatures A356 displayed higher ductility than at 120ºC, and for similar ΔK, the fatigue cracks growth rate was higher than that at 120ºC. However, in light of the deformation mechanism map presented by Roesler et al. (2007), for this temperature, two micro-mechanisms may cause plastic deformation: dislocation glide and dislocation creep [27]. Under low external stress and at low temperature, the material deforms elastically.

Similar to deformation at low temperatures, creep deformation can also occur by dislocation in metals. Nevertheless, there is one crucial difference: if a dislocation edge encounters an obstacle, e.g., a precipitate, it will need a minimum value of stress to overcome the obstacle at low temperatures. On the other hand, if a dislocation edge encounters an obstacle at higher temperatures, the dislocation can evade the obstacle by adding or emitting vacancies, thereby escaping its original slip plane. Diffusion creep starts at higher temperatures, being stronger under low stress than under dislocation creep because of its lower creep exponent. Due to the lower activation energy needed for grain boundary diffusion, this mechanism overrides bulk diffusion at low temperature. Since the creep exponent is the same in both cases, the two regions are separated by a vertical line.

Figure 13 shows the fracture surface of a specimen tested for fatigue at 280ºC under dwell cycles and load control. Crack propagation was mainly due to creep, with dimples appearing extensively at the crack front during dwell. Since a crack starts to propagate at relatively low ΔK, diffusion creep may have occurred in this case. As the crack grows, ΔK rises and crack growth quickly accelerate.

The fracture surface of a specimen tested for fatigue at 280ºC and under dwell cycles and displacement control, shown in Fig. 14, indicated a fatigue micro-mechanism of fracture similar to that under load control. However, in this case, it should be emphasized that stress relaxation takes place at the crack front, and ΔK decreases rapidly as a result while the crack propagation rate remains nearly constant. The fracture surface shows a mixture of dimples (creep) and mechanical fatigue features [28, 29].

Die-cast A356 Al alloy displayed a microstructure with rather coarse dendrites in a eutectic matrix, with mean SDAS of 25µm and intermetallic precipitates and porosities with a mean diameter size of 50µm and maximum size of 500µm.

Considering the isothermal fatigue tests under strain control, as the temperature rose from 120°C to 280, life reduced because of A356 becoming less resistant to plastic deformation and increasing number of defects. The stabilized hysteresis loop indicated a difference in ductility as well as the fact that during testing at 120°C, the material in question cyclically hardens while cyclically softening at 280°C.

In-phase thermo-mechanical test results indicated that the fatigue life of A356 is quite similar to that obtained from isothermal tests at 280°C and that the most severe damage took place at the highest temperature. Also, both hystereses are quite similar in shape, some disparities since the thermo-mechanical cycle are slower (0.003 Hz) than the isothermal one (0.1 Hz), which allows stress relaxation to occur, thereby reducing the maximum stress level.

Relaxation tests indicated two distinct behaviors, 240°C being the threshold temperature. At lower temperatures, A356 hardens cyclically while softening cyclically at higher temperatures. This behavior is similar to that observed during tests for isothermal fatigue. Given that the tests were carried out under strain control, there is stress relaxation during dwell, which reduces the elastic component of strain while concurrently increasing its plastic component. The higher the test temperature, the lower the amount of stress.

Results for fatigue crack growth show that temperature and waveform have a substantial impact on both crack growth rate and stress intensity threshold. About sinusoidal waveform (10 Hz), when the temperature increased, ΔKth decreased, and the crack propagation rate increased (for a constant K). However, the change in the rate of crack propagation by fatigue (da/dN) with temperature is quite similar, which indicates that the crack growth micro-mechanism does not change because of the high frequency used; it was due only to loss of mechanical strength instead.

For stress controlled testing at constant temperature (120 and 280°C), when the waveform changed from sinusoidal to 3-minute dwell at maximum load, the crack propagation rate increased by approximately 100 times, indicating a significant contribution from creep to crack propagation, with the amount of creep contribution dependent on temperature, i.e., the higher the temperature, the more significant the contribution from creep.

As regards strain controlled testing at constant temperature (280°C) and under 3-minute dwell at maximum load, the crack propagation rate (da/dN) as a function of ΔK became nearly constant because of stress relaxation.

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Fatigue Life Assessment of A356 Aluminum Alloy used for Engine Cylinder Head

Status:

Version 1

Abstract

Figures

Introduction

Material and Methods

Results and Discussions

Material

Mechanical Tests

Conclusions

References

Status:

Version 1