**9. Mechanical properties of Al-Cu-Mn-Zr alloy**

#### **9.1 Hardness and tensile properties**

Similar to other age hardening Al-Cu alloys, primary strengthening mechanism for peak-aged Al-Cu-Mn-Zr alloys is Orowan looping where matrix dislocations bow around the coarse θ<sup>0</sup> precipitates [52]. Apart from this, solid solution and grain boundary strengthening (by Hall–Petch mechanism) also contributes to the overall strength of the Al-Cu-Mn-Zr alloys. Analytical calculations, however, suggest that Orowan looping plus other strengthening mechanisms together are inadequate to account for experimentally measured yield strength of Al-Cu-Mn-Zr alloys [52]. This calls for the consideration of additional strengthening mechanisms e.g. stress-free transformation strain (SFTS). The formation of θ<sup>0</sup> precipitates in the α-Al matrix is usually associated with transformation strain fields which can interact and potentially restricts dislocation movement, thereby increasing the alloy strength further.

**Figures 3c–d** indicates that the room temperature tensile properties for peak aged Al-Cu-Mn-Zr alloy is inferior compared to the conventional peak-aged Al-Cu alloy. For example, the ultimate tensile strength(UTS) of Al-5Cu alloy is �490 MPa whereas it is �300 MPa for Al-Cu-Mn-Zr alloy at room temperature [48]. In addition, the yield strength of the later alloy is nearly half compared to the base Al-Cu alloy. However, after prolonged thermal exposure at 300°C, the trend reverses; Al-Cu-Mn-Zr alloy possess nearly twice the UTS and yield strength compared to the base Al-Cu alloy. Similarly, Al-5Cu-Mg alloy possess higher hardness than Al-Cu-Mn-Zr alloys at room temperature (**Figure 7**). With increase in pre-conditioning temperatures (heat treatment for 200 hours), non- Al-Cu-Mn-Zr alloys show drastic decrease in hardness around 200°C, while Al-Cu-Mn-Zr alloys can sustain the room temperature hardness without any significant degradation until 350°C.

Bahl. et al. [52] further showed that the hardness and yield strength of peak-aged Al-Cu-Mn-Zr alloy drop marginally during post-aging thermal exposure but remained

**Figure 7.** *Room temperature hardness for various conventional Al-Cu and Al-Cu-Mn-Zr alloys as a function of pre-conditioning temperatures [48].*

almost constant during prolonged thermal treatment up to 5000 hours. This accounts for a stable microstructure with almost constant precipitate volume fraction, thickness, diameter, aspect ratio, equivalent diameter, number density and interprecipitate spacing for Al-Cu-Mn-Zr alloys on extended thermal exposure. The ductility of Al-Cu-Mn-Zr alloys are further influenced from Cu content although it does not vary the yield strength and UTS much. For example, increasing the Cu content from 6 wt% to 9 wt% causes the fracture strain to reduce by 50% primarily due to the increased amount of brittle intermetallics at α-Al matrix grain boundaries [80].

Important to note that no comprehensive study is yet to report the results pertaining to full-scale tensile testing, especially the strain hardening response as well as the fracture characteristics of Al-Cu-Mn-Zr alloys whether at room or elevated temperatures. The earliest available work of Shyam et al. [48] showed certain true stress–strain curves for Al-Cu-Mn-Zr alloy from tensile tests carried out at room temperature and 300°C in comparison to regular Al-Cu alloy (**Figure 3**). The purpose of the tensile tests was however, to establish the superiority for the former alloy at elevated temperature. Important to note that the alloys (Al-Cu and Al-Cu-Mn-Zr alloys) were used in peak-aged condition for room temperature tensile tests and after pre-conditioning at 300°C for 200 hours for elevated temperature tests.

The true stress–strain curves from room temperature tensile tests suggest that peak-aged Al-Cu-Mn-Zr alloy possesses marginally higher strain hardening rate compared to the conventional Al-5Cu-Mg alloy at least in the initial part of the plastic regime. The hardening rates although do not vary much at the later part (below UTS) representing almost similar slopes for both alloys. During 300°C tensile tests, both Al-Cu-Mn-Zr and Al-5Cu-Mg alloys exhibit substantial strain softening, however at significantly differing rates; the rate of softening is greater for Al-Cu-Mn Zr alloy compared to the conventional Al-Cu-Mg alloy. The ductility for the former alloy is also always higher than the later alloy irrespective of the test temperature. The strain hardening response for precipitate hardened systems at room temperature is generally attributed to the isotropic hardening of α-Al matrix plus kinematic hardening due to dislocation pile up at the precipitate locations from continued Orowan looping [81–83]. The strain softening at elevated temperature can possibly be attributed to dynamic recovery of piled-up dislocations which reduces dislocation density at precipitate cites and increases ductility by delaying the final fracture.

*New-Age Al-Cu-Mn-Zr (ACMZ) Alloy for High Temperature-High Strength Applications… DOI: http://dx.doi.org/10.5772/intechopen.104533*

### **9.2 Creep response**

The excellent high temperature stability of Al-Cu-Mn-Zr alloys enhances their creep properties as well. Miligan et al. [84] conducted creep tests under different stress levels for various Al-Cu-Mn-Zr alloys with varying grain sizes at 300°C and compared their creep resistance with base Al-Cu alloy as well as Al-Sc alloy which is known for its excellent creep resistance. **Figure 8** represents the steady state creep strain rate as a function of applied stress for these alloys. At low stress level, stress exponents for Al-Cu-Mn-Zr alloys are close to unity signifying for diffusional creep being the dominant mechanism. On the other hand, dislocation creep is the mechanism for conventional Al-Cu alloy as identified from a higher stress exponent. The dislocation movement though α-Al grain interiors is difficult at low stress levels for Al-Cu-Mn-Zr alloys due to the enhanced thermal stability of θ<sup>0</sup> precipitates; rather grain boundary diffusion dominates at high temperature making diffusional Coble creep as the rate controlling mechanism. At higher stress levels however, the controlling mechanism switches to dislocation creep even for Al-Cu-Mn-Zr alloys since the grain boundary precipitates effectively slow down the movement of vacancies. This in turn restricts grain boundary diffusion as well as grain boundary sliding.

#### **9.3 High temperature deformation response**

One of the prime motivations for the development of Al-Cu-Mn-Zr alloy is to replace conventional cast Aluminum alloys (e.g. Al-Si-Cu based 319 alloy) for making light weight components in automotive engines [48, 85]. The Al-Cu-Ni based RR350 alloy with 0.2 wt% Mn and 0.17 wt% Zr, which can be considered as a variant of Al-Cu-Mn-Zr alloys, is also used as light weight and high temperature resistant alloys for high end automobile engine applications over the years [86]. Shower et al. [87] has compared the effect of microstructural stability on the high temperature deformation response of RR350 alloy vis-à-vis 319 alloy in as-cast

#### **Figure 8.**

*Creep curves showing the steady state creep strain rate as a function of applied stress for various Al-Cu-Mn-Zr alloys (RR350, Al-7Cu SG and Al-7Cu LG where SG and LG refers to small and large grains, respectively) plus base Al-Cu and Al-Sc alloys [84].*

**Figure 9.**

*(a) Comparison of flow stress variations with test temperatures for 319 and RR350 alloys at different strain rates, and (b) post-compression (at 300°C and 1 s*�*<sup>1</sup> strain rate) SEM micrograph of RR350 alloy showing shear band formation and bending of θ*<sup>0</sup> *precipitates within α-Al grains having <100> direction nearly parallel to the compression axis [87].*

condition by conducting isothermal hot compression tests at different temperature-true strain rate combinations. At all strain rates, compressive flow stress of 319 alloy is greater than that for RR350 alloy up to 200°C (**Figure 9a**). However, within 250–300°C, RR350 alloy possesses higher flow stress which can be attributed to the stability of strengthening θ<sup>0</sup> precipitates. In this temperature range, 319 alloy losses its flow stress by 40%. At lower strain rates (e.g. 10�<sup>4</sup> s �<sup>1</sup> & 10�<sup>3</sup> s �1 ), primary deformation mechanism for both alloys is strain hardening at room temperature, which changes to dynamic recovery and cross slip of dislocation at 250°C. Afterwards, dynamic recrystallization becomes predominant at 300°C while grain boundary sliding is the primary deformation mechanism at 350°C. In addition, RR350 alloy shows formation of shear bands as well as bending of θ<sup>0</sup> precipitates within α-Al grains having <100> direction nearly parallel to the compression axis in the microstructure of specimens deformed at 300°C (**Figure 9b)**.

### **10. Fatigue response**

The excellent high temperature stability of Al-Cu-Mn-Zr alloys also make them a prime candidate for fracture critical engineering applications where fatigue properties are crucial consideration. Bahl et al. [88] studied the effect of Cu concentration on the high temperature (250°C) low cycle fatigue (LCF) properties of Al-Cu-Mn-Zr alloys that led to a correlation between LCF life and monotonic tensile fracture strain. At low strain amplitude (0.1%), Al-Cu-Mn-Zr alloys with either 6 wt% or 9 wt% of Cu do not undergo failure even after 10<sup>5</sup> number of cycles. However, at higher strain amplitudes (0.2% and 0.3%), the alloys fail within these many cycles of testing. This suggests that the fatigue life of Al-Cu-Mn-Zr alloys decreases with increasing strain amplitude.

For the peak-aged Al-Cu-Mn-Zr alloys with varying Cu content that underwent a further thermal exposure at 250°C for 100 hours, initiation of fatigue cracks almost always occur from the surface pores rather than from coarse grain boundary precipitates [80]. The fracture surfaces correspondingly do not contain much of the traces of intermetallic precipitates. Finite element modeling (FEM) also indicates that stress concentration at the pores are higher compared to that at the grain

#### *New-Age Al-Cu-Mn-Zr (ACMZ) Alloy for High Temperature-High Strength Applications… DOI: http://dx.doi.org/10.5772/intechopen.104533*

boundary precipitates. The low strain amplitude in the fatigue testing possibly led to the pore assisted crack initiation since otherwise the cracking from grain boundary precipitates would require higher stress concentration and their decohesion from the matrix which is only possible at larger strains [88].

Furthermore, since the variation in Cu content only affects the volume fraction of grain boundary intermetallic precipitates, it does not influence the cracking and in turn, low cycle fatigue behavior of Al-Cu-Mn-Zr alloys [80]. The thermal stability of θ<sup>0</sup> precipitates also does not influence the fatigue property of these alloys since both crack initiation and propagation occur at a larger microstructural scale (from surface pores). It therefore appears that controlling the casting defects (predominantly shrinkage pores) is the most crucial factor to enhance the LCF life for Al-Cu-Mn-Zr alloys. Overall, these alloys exhibited moderate to excellent high temperature low cycle fatigue life making them suitable for components meant for elevated temperature applications.
