New-Age Al-Cu-Mn-Zr (ACMZ) Alloy for High Temperature-High Strength Applications: A Review

Abstract

One of the prime challenges with age hardened Al-Cu alloys is the strength degradation at high temperatures (above ∼250°C) due to the coarsening of strengthening θ′ precipitates and associated metastable θ′ → stable θ phase transformation. A recent discovery suggests that micro-alloying with Manganese (Mn) and Zirconium (Zr) can synergistically restrict θ′ precipitate coarsening, thereby rendering an excellent high temperature stability for Al-Cu-Mn-Zr (ACMZ) alloys. The θ′ precipitates are stabilized primarily from the reduction of interfacial energy by preferential solute segregation (Mn & Zr) at θ′ precipitate/α-Al matrix interfaces. The Al-Cu-Mn-Zr alloys thereby exhibit excellent high temperature hardness and tensile properties (yield and ultimate tensile strength) in addition to superior fatigue life and creep resistance. This newly developed Al-Cu-Mn-Zr alloys also showed excellent hot tearing resistance compared to the conventional cast Al-Cu alloys so much so that it meets the industrial standards as well. These alloys also have promising manufacturing possibility by additive route. Overall, Al-Cu-Mn-Zr alloys offer great potential for the automotive industry because of their unprecedented high temperature performance which should enable engineers to build light weight passenger vehicles leading to a safer and greener environment.

1. Introduction

Aluminum alloys have been one of the most prominent structural material system for many years now; this is also reflected in their global usage, they come only next to steel [1, 2]. Owing to their high specific strength, resistance to stress corrosion cracking, excellent fatigue resistance, workability and cost effectiveness [3], Aluminum alloys are one of the primary material choice for aerospace and automotive industries [2, 4]. In addition, Aluminum based composites were developed over the years to mitigate some of the limitations of Aluminum alloys and further facilitates their use for various engineering applications [5]. Apart from structural applications, Aluminum and its alloys are also employed for the electronics and electrical industries in abundance for their suitable combination of functional properties [6].

Pure Aluminum is characterized by low yield strength which is improved by many folds from different strengthening strategies e.g. by adding different alloying elements, thereby making the alloys suitable for structural applications. Depending upon the major alloying element/s, Aluminum alloys are classified in two important categories; some of the alloys can be strengthened by heat treatment (age hardening) while others by mechanical deformation (non-age hardenable) [7, 8]. Al-Cu alloys fall in the first category; they are strengthened by in situ precipitates introduced through appropriate thermal treatment (aging). In age hardening, these precipitates hinder the dislocation motion and increase the hardness or strength of the alloy [9].

Despite the beneficial attributes, age-hardening Aluminum alloys, especially Al-Cu alloys suffer from limited high temperature capability. At temperature above 200°C, the metastable strengthening precipitates (e.g. θ′) undergo rapid growth and coarsening and even transform to stable incoherent precipitates (metastable θ′ → stable θ precipitate) which are inefficient to restrict dislocation motion. This leads to a concurrent steep decrease in their high temperature load bearing capabilities. Due to such degradation, Al-Cu alloys are restricted for elevated temperature applications in spite of their light weight and high specific strength at room temperature.

In recent times, a new class of Al-Cu alloys is developed by suitable micro-alloying with Mn and Zr which possesses excellent stability for strengthening θ′ precipitates at and above 300°C. The Al-Cu-Mn-Zr, termed as ACMZ alloys, provide significant improvement in most of the elevated temperature mechanical properties including hardness, tensile strength, creep and fatigue resistance etc. The present chapter provides a detailed account of the development of Al-Cu-Mn-Zr alloys while highlighting the limitations of existing Al-Cu alloys in the first place. It alongside discusses about the underlying mechanisms responsible for their excellent high temperature stability and subsequently on various properties. The chapter finally access the possibility of industrial adaptation of this newly developed alloy system and expected industrial impacts in long run.

2. Precipitate formation and evolution in Al-Cu system

During age hardening of Al-Cu alloys, it is first heat-treated at temperatures where single-phase α-Al solid solution is formed; this process is known as solution treatment [7]. Afterwards, the alloy is rapidly quenched to room temperature which causes freezing of solute Cu atoms within the α-Al matrix, thus forming a super saturated solid solution (SSSS). The solute atoms afterwards can diffuse even at room temperature through this super-saturated α-Al matrix and form various Al-Cu precipitates; this process is known as natural aging. In this regard, Figure 1a represents the binary Al-Cu phase diagram along with the solvus lines for various metastable precipitate phases. However, in common engineering practice, the solutionized alloy is heat treated at certain elevated temperature, usually within the two-phase region to produce various metastable precipitates depending on the heat treatment time and temperature (artificial aging). At sufficiently low aging temperature or during natural aging, Cu solute cluster first form from the quenched-in vacancies within the supersaturated α-Al matrix. These solute clusters then arrange in a single layer of Cu atoms known as $\mathrm{GP}$ -I zones (Guinier Preston zone) along a plane parallel to the (001) plane of the α-Al matrix.

Figure 1.

Schematics showing (a) section of Al-Cu binary phase diagram, (b) crystal structure of parent α-Al matrix and various metastable (${\mathrm{\theta }}^{\prime \prime }$ and ${\mathrm{\theta }}^{\prime }$) and stable ($\mathrm{\theta })$ precipitates in Al-Cu alloy system, (c) morphology and different interfaces for metastable ${\mathrm{\theta }}^{\prime }$ precipitate.

When the aging process continues above the GP zone solvus line (Figure 1a), ${\mathrm{\theta }}^{\prime \prime }$ metastable precipitates (also known as $\mathrm{\text{GP}}$-II zones) form by diffusion of Cu atoms from the solute cluster zones ($\mathrm{GP}$ -I zone). ${\mathrm{\theta }}^{\prime \prime }$ are characterized by the chemical formula Al₃Cu and ordered tetragonal structure ($a=b=4.04Å,c=7.68Å$) having two layers of Cu atoms separated by three layers of Al atoms (Figure 1b) [10]. It is completely coherent with α-Al matrix from all sides although misfit strain develops along ${\left(100\right)}_{\mathit{Al}}$ and ${\left(010\right)}_{\mathit{Al}}$ planes. Metastable ${\mathrm{\theta }}^{\prime }$ precipitates next form from ${\mathrm{\theta }}^{\prime \prime }$ precipitates on further aging. ${\mathrm{\theta }}^{\prime }$ has a stoichiometric formula of Al₂Cu with body centred tetragonal structure ($a=b=4.04Å,c=5.80Å$, space group I4/mmm) and carries 3 layers of Al atom and 2 layers of Cu atoms in an ordered arrangement (Figure 1b) [10, 11].

In the ${\mathrm{\theta }}^{\prime \prime }$→${\mathrm{\theta }}^{\prime }$ transformation, the later precipitate starts nucleating at the pre-existing ${\mathrm{\theta }}^{\prime \prime }$ precipitates via continued diffusion of Cu atom and continues to grow until the entire ${\mathrm{\theta }}^{\prime \prime }$ precipitate transforms into ${\mathrm{\theta }}^{\prime }$. In this transformation, ${\mathrm{\theta }}^{\prime }$ holds the same orientation relationship with α-Al matrix as does ${\mathrm{\theta }}^{\prime \prime }$ precipitate i.e. ${\left(001\right)}_{{\mathrm{\theta }}^{\prime }}\Vert {\left(001\right)}_{\mathit{Al}}$ and ${\left[100\right]}_{{\mathrm{\theta }}^{\prime }}\Vert {\left[100\right]}_{\mathit{Al}}$ [12]. It grows as a plate within the α-Al matrix parallel to the ${\left(010\right)}_{\mathit{Al}}$ and remains coherent along this plane. However, along the ${\left(100\right)}_{\mathit{Al}}$ and ${\left(010\right)}_{\mathit{Al}}$ planes, ${\mathrm{\theta }}^{\prime }$ becomes incoherent or complex semi-coherent with α-Al matrix. The semi-coherent side of the plate shaped ${\mathrm{\theta }}^{\prime }$ precipitate is not completely circular; rather it forms as octagon with facets along ${\left[100\right]}_{{\mathrm{\theta }}^{\prime }}$ and ${\left[110\right]}_{{\mathrm{\theta }}^{\prime }}$ directions (Figure 1c). The ${\left\{110\right\}}_{{\mathrm{\theta }}^{\prime }}$ interface edges further act as a solute gateway from where Cu atoms diffuse and coarsen ${\mathrm{\theta }}^{\prime }$ precipitates on prolonged thermal exposure. The atomic arrangement as well as ledge dislocations on ${\left\{110\right\}}_{{\mathrm{\theta }}^{\prime }}$ semi-coherent interfaces further assist in the accommodation of Cu atoms [13].

The equilibrium precipitate in the Al-Cu system is tetragonal $\mathrm{\theta }$ ($a=b=6.07Å,c=4.87Å$, space group I4/mcm) having chemical formula Al₂Cu (Figure 1b) [10]. It is incoherent with the α-Al matrix along all sides. $\mathrm{\theta }$ forms from ${\mathrm{\theta }}^{\prime }$ precipitate via mechanism similar to ${\mathrm{\theta }}^{\prime \prime }$ → ${\mathrm{\theta }}^{\prime }$ transformation. The habit plane of $\mathrm{\theta }$ also remains parallel to (100) α-Al matrix planes.

The entire precipitation process during aging of Al-Cu alloy therefore can be summed up as: $\text{Solute clusters}\to \mathrm{GP}\text{zones}\to \mathrm{GP}\mathrm{II}\text{zones}\left({\mathrm{\theta }}^{\prime \prime }\right)\to {\mathrm{\theta }}^{\prime }\to \mathrm{\theta }.$

3. Strengthening mechanisms in Al-Cu alloy system

The operating strengthening mechanism/s in Al-Cu alloy system differs as a function of precipitate type, mainly with their size and coherency with α-Al matrix. When the precipitates are small and coherent with α-Al matrix (e.g. $\mathrm{GP}$-I or ${\mathrm{\theta }}^{\prime \prime }$), it is energetically easy for the dislocations to shear through them. This usually occurs via one or a combination of mechanisms like chemical strengthening, stacking-fault strengthening, modulus strengthening, order strengthening, coherency strengthening etc. [14]. Apart from chemical strengthening, the increment in critical resolved shear stress (CRSS) varies with the size of the precipitates as ∼ r^1/2 where r is the precipitate radius. In case of chemical strengthening, CRSS increment is inversely proportional to r.

When the precipitates (usually ${\mathrm{\theta }}^{\prime }$) are large in size and possess semi-coherent or incoherent interfaces with α-Al matrix, the dislocation line bulges within the inter-precipitate region rather than shearing through the precipitate until they meet and move forward while leaving a dislocation loop behind (Figure 2a) [16]. This process is known as Orowan looping. The increase in CRSS (for spherical precipitates) due to Orowan looping is given by:

Figure 2.

Schematic showing (a) the mechanism of Orowan looping, (b) CRSS increment as a function of precipitate radius for particle shearing and Orowan looping mechanisms, and (c) isothermal aging curves (hardness vs. aging time) for various commercial Al alloys e.g. Al-Cu (206), Al-Si-Cu (319) and Al-Si (356, A356, A356 + 0.5Cu) [15].

$$∆\tau ={\left(\frac{3}{2\pi }\right)}^{\frac{1}{2}}.\frac{\mathit{Gb}}{r}.f^{\frac{1}{2}}$$E1

where, $∆\tau$ is the increase in CRSS due to Orowan strengthening, G is the shear modulus of the α-Al matrix and $f$is the precipitate volume fraction. The CRSS increment thus varies with 1/r. Furthermore, for plate shaped ${\mathrm{\theta }}^{\prime }$ precipitate forming on {100}_Al habit plane having diameter $D$ and thickness $t,$ CRSS increment, $∆\tau$ due to Orowan looping is given by [17]:

$$∆\tau =\left(\frac{\mathit{Gb}}{2\pi \sqrt{1-\nu }}\right)\left(\frac{1}{0.931\sqrt{\frac{0.306\mathit{\pi DT}}{f}-\frac{\mathit{\pi D}}{t}-1.061T}}\right)\left(\mathit{ln}\frac{1.225t}{r_0}\right)$$E2

where, $\nu$ is the Poisson’s ratio of α-Al matrix and $r_0$ is the radius of the dislocation core.

The CRSS increment with respect to the precipitate radius due to either particle shearing or Orowan looping is schematically represented in Figure 2b [12]. It seems that a critical radius exists for the strengthening precipitate below which particle shearing is preferred. When the precipitate grows beyond this critical radius, the dislocations prefer to bow around the precipitate rather than shearing it.

Considering finely dispersed coherent ${\mathrm{\theta }}^{\prime \prime }$ precipitates within α-Al matrix, they cause strength increment by one or a combination of mechanisms mentioned above. As these coherent precipitates grow, a corresponding increase occurs in the alloy strength since CRSS increment is proportional to r^1/2 for most of the strengthening mechanisms. At the critical radius, the strength is highest, also signifying for ${\mathrm{\theta }}^{\prime \prime }\to {\mathrm{\theta }}^{\prime }$ transformation. Afterwards, dislocation bowing around large and semi-coherent ${\mathrm{\theta }}^{\prime }$ precipitates is the main strengthening process. On further coarsening, the number density continuously decreases for ${\mathrm{\theta }}^{\prime }$ precipitates so that the alloy strength correspondingly decreases since CRSS increment is proportional to 1/r in Orowan looping.

4. Age hardening and aging curve

At the beginning of aging treatment, fine and uniform precipitation of $\mathrm{GP}$ -I and/or ${\mathrm{\theta }}^{\prime \prime }$ occurs within the super saturated α-Al matrix. These precipitates hinder the dislocation motion and increases the strength (or hardness) compared to the solutionized and quenched alloy (Figure 2c) [7, 8]. The number density of the precipitates is less and inter-precipitate distance is correspondingly high at this condition so that the alloy strength/hardness is marginally increased; the alloy is said to be in under-aged condition. Shearing of coherent $\mathrm{GP}$-I and ${\mathrm{\theta }}^{\prime \prime }$ precipitates is the primary strengthening mechanism here. As the aging time increases, $\mathrm{GP}$ -I zones continuously transform to ${\mathrm{\theta }}^{\prime \prime }$, thereby increasing their number density in the α-Al matrix. Some of them even transform to other metastable precipitates e.g. ${\mathrm{\theta }}^{\prime }$. Together, the strength/hardness increases further with continued aging due to shearing of coherent and semi-coherent precipitates [18].

With increase in aging time, aspect ratio and number density of ${\mathrm{\theta }}^{\prime }$ precipitates continuously increase because of which precipitate shearing becomes more and more difficult; the strength/hardness of the alloy also keeps increasing gradually at this stage [19]. At a critical precipitate size, the strength (or hardness) of the alloy reaches a peak denoting the peak-aged condition. The ${\mathrm{\theta }}^{\prime }$ precipitate size, aspect ratio, inter-precipitate distance etc. are now optimum for maximum hindrance towards dislocation motion due to shearing through the precipitates.

The ${\mathrm{\theta }}^{\prime }$ precipitates continuously coarsen with increasing aging time and even start transforming to stable θ precipitate under prolonged aging [12]. The energy required to shear ${\mathrm{\theta }}^{\prime }$ precipitates becomes quite high at this stage so that rather than shearing, dislocations prefer to bow around them (Orowan looping) [8]. The strength/hardness of the alloy thereafter decreases with increasing aging time leading to over-aged condition when metastable ${\mathrm{\theta }}^{\prime }$ continue transforming to stable θ precipitate. Figure 2c represents typical aging curves (hardness vs aging time) for several commercial age-hardening Al alloys e.g. Al-Cu (206), Al-Si-Cu (319) and Al-Si (A356, A356 + 0.5Cu). In these aging curves, hardness of the alloys initially increase (under-aged condition), reaches the maximum at peak-aged condition and decreases again with subsequent aging leading to the over-aged condition [12].

Furthermore, the ductility (usually expressed in terms of elongation to fracture) of age-hardening Aluminum alloys also varies with aging time in accordance to the size, morphology and coherency of the strengthening precipitates [20, 21]. When the precipitates are small and coherent and their number density is low, dislocations can move past easily through them leading to maximum contribution from strain hardening that delays fracture. Correspondingly, under-aged alloy shows maximum ductility. On the other hand, semi-coherent and large strengthening precipitates at the peak aged condition renders maximum hindrance to the dislocation motion due to Orowan looping resulting in significant pile up at the precipitate sites. The ductility for the peak-aged alloy is also correspondingly minimum. With over-aging, some ductility is restored since the coarse, incoherent precipitates are generally not suitable to hinder dislocation motion and little pile up results around them. Overall, the ductility vs. aging time variation follows a reverse trend to the strength/hardness vs. aging time curves.

5. Challenge for high temperature stability of Al-Cu alloy: precipitate coarsening

One of the major hindrances for widespread use of Al-Cu alloys, especially in the automobile and aerospace sectors, is their poor high temperature stability associated with rapid decrease in load bearing capacity above ∼250°C [22]. This strength degradation is because of the rapid coarsening of θ′ precipitates from thickening along the broad facets, thereby resulting in a drastic decrease in their aspect ratio. At high temperature, the rate of diffusion for solute Cu atoms increases so that they segregate at the coherent and semi-coherent interfaces of ${\mathrm{\theta }}^{\prime }$ precipitates [23]. The broad facet (i.e. the coherent interface) thickens by ledge formation while the semi-coherent interfaces having high interfacial energy grow from accommodation of Cu atoms along {100}_θ′, {010}_θ′, {110}_θ′ interfaces [24]. The thermodynamic driving force for coarsening of metastable ${\mathrm{\theta }}^{\prime }$ precipitates is the reduction of interfacial energy [13]. It should be noted however that the increase in diameter of ${\mathrm{\theta }}^{\prime }$ precipitate by the growth of semi-coherent interface does not essentially reduce the interfacial energy; rather the thickening of the broad facet i.e. growth of the coherent interface is responsible for the decrease in interfacial energy during coarsening.

At long thermal exposure, ${\mathrm{\theta }}^{\prime }$ precipitates ultimately transform to stable (equilibrium) θ precipitates, which are completely incoherent with the parent α-Al matrix. Further thermal treatment leads to the growth of larger θ precipitates in expense of the smaller ones. This corresponds to an increase in the inter-precipitate distance so that coarse θ precipitates are no longer effective in restricting the motion of matrix dislocations; hence, the strength of the alloys decreases drastically [9]. It therefore appears imperative to stabilize strengthening metastable ${\mathrm{\theta }}^{\prime }$ precipitates against coarsening for successful high temperature application (beyond ∼250°C) of Al-Cu alloys.

6. Attempts towards the development of high temperature Aluminum alloys

Over the years, alloying pure Aluminum with various elements showed good promises for elevated temperature applications. Such elements included rare earths (e.g. Erbium, Ytterbium, Scandium etc.) as well as Zirconium, Silicon etc. [25, 26, 27]. For these alloys, formation of coherent precipitates with cubic L1₂ crystal structure and reduced interfacial energy is the key for their high temperature stability against precipitate coarsening upon thermal exposure [26]. The other unique exploration is the formation of core-shell structure for the strengthening precipitates which also provides excellent coarsening resistance through the minimization of interfacial energy. For example, addition of 0.06 at% Zr or 0.03 at% Er in Aluminum individually form ordered L1₂ Al₃Zr or Al₃Er precipitates and show moderate coarsening resistance at elevated temperature [25]. Simultaneous addition of Zr and Er in similar quantity however, leads to Al₃(Er, Zr) precipitates with unique core shell structure, which made them coarsening resistant up to 400°C for 750 hours. The difference in diffusivity between Zr and Er was held responsible for formation of such core-shell structure; while Er having higher diffusivity forms the primary precipitate with Aluminum, slower diffusing Zr segregates later at the interfaces of these primary precipitates resulting in the core-shell structure. Similarly, addition of Sc to Al-Zr-Sc-Er alloy (concentrations of both Sc and Er are 0.06 at%) leads to a dual shell layer of Zr and Sc according to their respective diffusivity in Al matrix over the Al₃Er core precipitate [26].

Furthermore, excellent creep resistance was observed for Al-0.1 at% Zr and Al-0.1 at% Zr-0.1 at% Ti alloy systems at 300°C, 350°C and 400°C, which is attributed to the high temperature stability of Al₃Zr precipitates [28]. Out of the two alloys, ternary Al-Zr-Ti alloy showed comparatively lower creep resistance than binary Al-Zr alloy due to the lower lattice parameter mismatch between Al₃(Zr_1-xTi_x) core-shell precipitates with the parent α-Al matrix. The addition of Yb similarly resulted in excellent thermal stability for Al-0.9 at% Zr- 1.73 at% Yb alloys having Al₃(Zr,Yb) precipitates up to 400–425°C [29].

For Al-Si system, Al-Si-Cu-Mg alloys are traditionally used for making high temperature pistons for automobile engines [30, 31]. These alloys show satisfactory microstructural stability as well as fatigue resistance at high temperatures which are essential requirements for automotive applications. A viable route for further improving their high temperature performance is by addition of transition metals that forms thermally stable intermetallic precipitates. For example, controlled Zr addition (up to 0.11 wt%) increases the ultimate tensile strength (UTS) of Al-Si-Zr piston alloys by 3.8% at 350°C due to the alteration in the morphology of strengthening ZrAlSi precipitates from flake to block shape [32, 33]. However, increase in Zr content up to 0.46 wt% resulted in a decrease in UTS by 5%. Similarly, A356 alloy (Al-7Si-0.4 Mg) modified with 0.25 wt% Er and nominal amount of Zr (0 to 0.6 wt%) showed improved high temperature mechanical properties [34]. With increase in Zr content up to 0.59 wt%, both hardness and tensile strength increases at room and elevated temperatures due to the formation of Al₃(Er,Zr) precipitates.

Hypoeutectic Al-7 wt%Si-1wt%Cu-0.5 wt% Mg alloys also shows excellent retention of hardness and tensile strength up to 240–260°C when micro-alloyed with 0.15 wt% Zr, 0.28 wt% V and 0.18 wt% Ti [35]. Further exposure to 475°C upto 128 hours led to additional improvement in hardness which can be attributed to the accelerated precipitation of Al₃(Zr,V,Ti) and Q′ precipitates. Similarly, addition of minor Ti (0.22 wt%), Zr (0.39 and 0.19 wt%) and Ni (0.46–0.21 wt%) to commercial 354 alloy showed improvement in tensile properties up to 300°C compared to the base alloy [36]. In both cases, micro-alloying elements synergistically result in unique and complex precipitate formation which improved the high temperature stability of the corresponding alloys. For hypereutectic Al-Si alloys, Ni addition up to 1–4 wt% to Al-12 wt%Si-0.9 wt% Cu-0.8 wt% Mg alloy resulted in retention of room temperature mechanical properties, including creep resistance up to 250°C due to the formation of thermally stable Al₃Ni precipitates [37]. In addition to primary and eutectic Si, incorporation of 1 wt% ZnO nanoparticles (particle size ∼40 nm) also enhance the high temperature tensile strength and elongation for Al-20 wt% Si alloys [38].

In case of age hardening Al-Cu system, several attempts were made in the past to increase their high temperature stability by adopting various strategies. Lin et al. [39] studied the effect of Ni addition (0.5–1.5 wt%) on the elevated temperature mechanical properties of squeeze cast Al-Cu-Mn-Fe alloys. At 300°C, the amount of thermally stable precipitates (e.g. Al₉FeNi, Al₃CuNi and Al₂₀Cu₂Mn₃) increases with increasing Ni content which enhances the elevated temperature mechanical properties of the base alloy. Addition of La in Al-Cu alloy similarly results in the formation of Al₁₁La₃ precipitates leading to a better high temperature mechanical properties with 0.3 wt% La being the optimized concentration [40]. The addition of 1.6–2.0 wt% Li also shows excellent mechanical properties for AA2099 (Al-Cu-Li) alloys at high temperature, primarily due to the enhanced thermal stability of T₁ (Al₂CuLi) precipitates compared to other possible strengthening precipitates like θ′ and S (Al₂CuMg) [41]. At higher temperature, T₁ precipitates coarsen instead of dissolving unlike θ′ or S. In addition, AA2219 alloy possesses improved high temperature performance when micro-alloyed with 0.8 wt% Sc, 0.45 wt% Mg and 0.2 wt% Zr from grain refinement and simultaneous precipitation of Al₃Sc, Al₃Zr and Ω precipitates along with other common strengthening precipitates like θ′ and θ″ [42].

Another viable strategy of increasing the thermal stability of Al-Cu alloys is by micro-alloying with various secondary elements for stabilization of strengthening metastable ${\mathrm{\theta }}^{\prime }$ precipitates. For example, 0.18 at% Sc addition forms thermally stable Al₃Sc precipitate wherein Sc atoms tend to segregate at the α-Al matrix/θ′ precipitate interfaces [43]. Such co-stabilization of two different precipitates renders excellent high temperature property for Al-Cu alloys with Sc addition [44]. Micro alloying with Zr provides similar effects as Sc; Zr forms stable Al₃Zr precipitates having L1₂ structure which can act as a heterogeneous nucleation site for θ′′ precipitates, thus generating a finer scale microstructure for Al-Cu-Zr alloys [45]. The resultant alloy shows excellent thermal stability up to 250–300°C. In Al-Cu-Mg alloys, 0.09–0.13 wt% Mg is reported to accelerate the formation of θ′ precipitates and the resultant alloy exhibits excellent thermal stability at 300°C for 1000 hours [46]. Similarly, addition of 0.45 wt% Ce provides heterogeneous nucleation sites for precipitation of Ω precipitates and enhances the thermal stability of Al-5.3% Cu-0.8% Mg-0.6% Ag (wt.%) alloys by retarding the diffusion of Cu [47].

Overall, there have been numerous efforts in the past to design high temperature Aluminum alloys from different binary systems (Al-Cu, Al-Si etc.), primarily by micro-alloying with various elements. However, most of these attempts showed certain shortcomings. The working temperature of the resultant ternary or quternary alloys could not be increased above 300°C under prolonged exposure. Also, use of exotic elements like rare earth additions hindered their industrial acceptance and commercial viability. Hence, the demand of cost-effective Aluminum alloys for high temperature applications has only increased over the years without much of a success.

7. Development of Al-Cu-Mn-Zr alloy

As mentioned before, age hardening Al-Cu alloys faces significant precipitate coarsening, which restricts their use for high temperature applications [48]. Numerous attempts have made over the years to increase the operational temperature for Al-Cu alloys; the most successful approach was by trace addition of various elements like Sc and Zr [49]. Micro-alloying improves the high temperature stability in two distinct ways:

1. Micro alloying elements can provide heterogeneous nucleation sites for primary strengthening precipitates (e.g. θ′′ and θ′) so that relatively finer precipitates with narrow size distribution is obtained in the room temperature microstructure [50]. For instance, when Al-Cu alloys are micro-alloyed with Sn, the semi-coherent interface of metastable θ′ precipitate nucleates from Sn particle [51].

2. Micro alloying elements can segregate at the high energy mobile semi-coherent precipitate/matrix interfaces and at times, at the less-mobile coherent interfaces [52]. Such segregation eventually results in the reduction of energy for these interfaces, making them difficult to grow. The semi-coherent interfaces is usually more effected by such segregation [53, 54]. In recent times, stabilization of semi-coherent and coherent interfaces of strengthening θ′ precipitates yield a new series of high temperature Al-Cu alloys with unprecedented thermal stability up to 350°C and beyond [48]. These alloys contain micro-alloying addition of Manganese and Zirconium and designated as Al-Cu-Mn-Zr or ACMZ alloys.

The classical approach for developing any new alloy system relies on the age-old trial and error method which has serious drawbacks, primarily considering the resource constrains and added cost from industrial standpoints [55, 56]. A more state of the art strategy of alloy designing is by using integrated computational materials engineering (ICME) approach [57, 58]; a successful example of this is realized in the development of Al-Cu-Mn-Zr alloys. The key components of ICME approach for the development of this alloy system are: (a) thermodynamic and kinetic approximations for stability of precipitates against growth controlling mechanism/s, (b) appropriate modeling for assessment of thermo-physical and thermo-mechanical properties from existing phases, (c) simulation and model/s to predict defect formation during casting processes, (d) models for prediction of microstructure during casting and other thermo-mechanical processing operations, (e) models for property prediction from microstructure and defect structure evolution, and (f) models for manufacturing of components at in-service conditions [59, 60].

These above mentioned steps were followed in the development of Al-Cu-Mn-Zr alloys which was primarily aimed to replace traditional Al-Si and Al-Si-Cu alloys in automotive applications (e.g. cylinder heads in passenger vehicle engines) [59]. Firstly, thermo-physical and thermo-mechanical properties for casting process simulation were obtained from thermodynamic databases. The simulation of casting process was conducted to estimate the casting defects and as-cast microstructure. Afterwards, thermodynamic models were employed for optimization of heat treatment cycle in terms of desired precipitation sequence, precipitate growth etc. This helped to estimate the spatial variation of thermo-physical properties over the component scale as well. Finally, component level properties e.g. residual stress, fatigue performance etc. were evaluated in order to estimate the in-service performance of the alloys considering the above mentioned parameters [61].

8. High temperature stability of Al-Cu-Mn-Zr alloy

The primary mechanism for high temperature stability of Al-Cu-Mn-Zr alloys is related to the segregation of micro-alloying solute atoms (Mn and Zr) at θ′ precipitate/α-Al matrix interfaces [48, 52]. Although it seems fairly straight forward in the first go, the interface stabilization process exhibits extreme intricacies throughout the entire precipitation sequence. As denoted earlier, the primary strengthening precipitate in Al-Cu system is θ′, which has a plate shaped morphology where the broad facets are coherent with parent α-Al matrix (Figure 1c) [12]. The rim of the precipitates, on the other hand, are semi-coherent and have a higher interfacial energy compared to their coherent counterparts which makes them highly mobile and prone to coarsening [48].

On exposure to high temperature for an extended duration, θ′ precipitates coarsen due to the enhanced diffusion of solute Cu atoms [8]. Segregation of Mn and Zr atoms at the θ′ precipitate/α-Al matrix interfaces prohibits Cu diffusion and further coarsening at elevated temperature [48]. The main driving force behind the solute segregation is the reduction of precipitate/matrix interfacial energy, especially for the semi-coherent interfaces. In addition, several other mechanisms like solute drag, ledge poisoning etc. also helps in the stabilization of θ′ precipitates [52, 62]. These mechanisms are explained individually below.

8.1 Segregation of micro-alloying elements

In the earliest report on Al-Cu-Mn-Zr alloys, Shyam et al. [48] compared two cast Al-Cu-Mn-Zr alloys having nominal compositions Al-5Cu-1.5Ni-0.2Mn-0.17Zr and Al-6.4Cu-0.19Mn-0.13Zr (in wt%) with conventional Al-Cu (206) and Al-Si-Cu (319) i.e. with non- Al-Cu-Mn-Zr alloys containing negligible concentration of Zr (Figure 3). At room temperature, base Al-Cu and Al-Si-Cu alloys exhibit superior mechanical response (higher yield strength and ductility) than Al-Cu-Mn-Zr alloys. However, the trend completely reverses after treating the alloys at higher temperature (300°C) for 200 hours; Al-Cu-Mn-Zr alloys now represent superordinate mechanical response than either Al-Cu or Al-Si-Cu alloys. Microstructural examinations reveal that θ′ precipitates significantly coarsen and transform to thermodynamically stable θ precipitates for base Al-Cu or Al-Si-Cu (i.e. non Al-Cu-Mn-Zr) alloys because of which their mechanical properties degraded after thermal treatment. On the other hand, θ′ precipitates retain their morphology and aspect ratio on high temperature heat treatment in case of Al-Cu-Mn-Zr alloys.

Figure 3.

(a) and (b) showing the microstructures of Al-Cu-Mn-Zr alloy in peak aged condition and post 300°C thermal exposure for 200 hours, respectively; (c) and (d) represents true-stress-true strain curves for these alloys from tensile tests carried out at room temperature and 300°C, respectively; (e) and (f) showing the microstructures for conventional Al-Si-Cu alloy under similar conditions [48].

Bahl et al. [52] further studied the aging kinetics and thermal stability of Al-Cu-Mn-Zr alloys and showed that they retained their room temperature mechanical properties even after exposure at 300°C for 5000 hours. During such prolonged thermal treatment, θ′ precipitates suffer limited decrease in number density up to 200 hours. No further significant decrease was observed, and the peak-aged microstructure remains fairly stable up to 5000 hours.

The stability of θ′ precipitates in Al-Cu-Mn-Zr alloy was examined using atom probe tomography (APT) characterization which are shown in Figures 4a and b (side view and top view, respectively) [48, 62]. Figure 4c represents corresponding composition profile which indicates segregation of Cu, Mn, Zr and Si at the coherent and semi-coherent θ′ precipitate/α-Al matrix interfaces for Al-Cu-Mn-Zr alloys after prolonged (200 hours) thermal exposure at 300°C. As it appears, Mn tends to segregate both at the coherent and semi-coherent interfaces of θ′ precipitates with the segregation tendency being larger at the later interfaces. Zr, on the other hand, segregates more on the corner rim of the coherent/semi-coherent interfaces, although certain extent of Zr segregation also occurs on these interfaces. Silicon have similar segregation profile as Mn; it can in fact influences the solute segregation at these interfaces to a much greater extent as discussed later [63].

Figure 4.

(a) and (b) APT compositional maps (iso-concentration surfaces) representing side and top views, respectively of APT needle for Al-Cu-Mn-Ni-Zr alloy pre-conditioned at 300°C for 200 hours, (c) and (d) showing 2D contour plots for Cu, Si, Zr and Mn atoms on the cross-sectional planes of θ′ precipitate along <110> direction from Al-5Cu-Ni-Mn-Zr and Al-7Cu-Mn-Zr alloys, respectively after pre-conditioning at 300°C for 200 hours [48].

In order to understand the individual and synergistic effect of Mn and Zr micro-alloying on the thermal stability of Al-Cu-Mn-Zr alloy, a consolidated study was carried out by Poplawsky et al. [62] on several model alloys (e.g. Al-Cu-Mn, ACM and Al-Cu-Zr, ACZ etc.) in addition to the base Al-Cu-Mn-Zr alloy. The Al-Cu-Mn alloys retain their room temperature mechanical strength after exposure at 300°C for 200 hours whereas Al-Cu-Zr alloys could sustain their stability only up to 200°C. For comparison, Al-Cu-Mn-Zr alloys are stable up to 350°C. The trend in Mn segregation for Al-Cu-Mn-Zr alloy in this case is similar to that observed earlier by Shyam et al. [48] up to 300°C. Larger Mn segregation occurs at semi-coherent interfaces while minor segregation at the coherent interfaces.

After thermal exposure at 350°C, Mn segregation at semi-coherent interfaces becomes insignificant, which aggravates the mechanical degradation of Al-Cu-Mn alloys at this temperature range [62]. At 350°C, Mn tends to diffuse within the bulk of θ′ precipitates, thereby causing even lesser segregation at the semi-coherent interfaces. Zr, on the other hand, retains their segregation profile at the coherent interfaces up to 200°C for Al-Cu-Zr alloys. The Zr segregation profile is also similar in nature to that observed previously for Al-Cu-Mn-Zr alloys by Shyam et al. [48].

8.2 Diffusional perspective

Solid state diffusion is one of the key component for evolution of precipitate structure and morphology during the course of thermal exposure [64]. In the corresponding binary systems, self-diffusion coefficient of Cu is much higher than that for Mn so that diffusion of Cu atoms continue to coarsen ${\mathrm{\theta }}^{\prime }$ precipitates unless interfacial segregation resists [63]. Furthermore, self-diffusion coefficient of Zr in Al is almost 10 times lower than that for Mn in Al [65, 66]. As a result, Mn atoms diffuse much faster to the coherent and semi-coherent interfaces of ${\mathrm{\theta }}^{\prime }$ precipitates than Zr atoms during the initial thermal exposure for peak-aged Al-Cu-Mn-Zr alloys. Mn segregation thereafter pins both the interfaces and restricts the flux of Cu atoms from coarsening the ${\mathrm{\theta }}^{\prime }$ precipitates on further heat treatment. The slow diffusing Zr atoms, on the other hand, tend to segregate mostly at the corner of the coherent and semi-coherent interfaces at a later stage of heat treatment and contributes to their stabilization only during prolonged thermal exposure.

Due to this sequence of segregation (initial segregation of Mn followed by Zr segregation on prolonged thermal exposure), ${\mathrm{\theta }}^{\prime }$ precipitates are stable only up to 300°C in Al-Cu-Mn (ACM) alloys where Mn alone is the micro-alloying element. In the absence of Zr, Mn segregation is insufficient to restrict coarsening at 350°C since they diffuse into the bulk of the precipitates rather than segregating at the interface at higher temperature or up to prolonged thermal exposure [62]. On the other hand, Zr being a slowly diffusing element, requires longer duration or higher temperature to segregate at the interface of α-Al/${\mathrm{\theta }}^{\prime }$ precipitates. However, sufficient coarsening of ${\mathrm{\theta }}^{\prime }$ precipitates may have already occurred or they may even transform to the stable θ precipitate by the time Zr stabilizes the interfaces on prolonged thermal exposure. In addition, Zr segregation preferentially takes places at the coherent interfaces which has lower mobility compared to the semi-coherent ones. Hence Zr segregation alone is least efficient to stabilize θ′ precipitate and the stability of Al-Cu-Zr (ACZ) alloys is limited only up to 200°C. In case of Al-Cu-Mn-Zr (ACMZ) alloys, Mn atoms initially stabilize both coherent and semi-coherent interfaces. The slow diffusing Zr atoms thereafter segregate at the coherent interfaces and provides further stabilization of ${\mathrm{\theta }}^{\prime }$ precipitates. This sequential segregation of Mn and Zr synergistically provides stability for Al-Cu-Mn-Zr alloys up to 350°C and beyond for a prolonged duration [48].

8.3 Effect of precipitate size and interparticle spacing

The diffusion aided coarsening of θ′ precipitates can be best described using the classic Lifshitz-Slyozov-Wagner (LSW) theory [67] where the rate of coarsening depends on the corresponding mass transfer mechanism (lattice diffusion, interface atomic mobility, grain boundary diffusion, pipe diffusion through dislocation cores etc.). The governing equation in LSW theory is given as [68]:

$${\overline{r}}^3-{{\overline{r}}_0}^3=k.t$$E3

where, $\overline{r}$ & ${\overline{r}}_0$ are the mean precipitate radius at time $t=t$ and $t=0$, respectively and $k$ is a constant. Partial derivation of Eq. (3) with respect to $t$ gives:

$$\frac{\partial \overline{r}}{\partial t}=\frac{k}{3{\left(k.t+{{\overline{r}}_0}^3\right)}^{\frac{2}{3}}}$$E4

Since $\frac{\partial \overline{r}}{\partial t}$ varies with the initial precipitate radius as $\sim \frac{1}{{{\overline{r}}_0}^2}$, the smaller the initial precipitate radius, the higher will be the rate of its coarsening. In case of peak-aged Al-Cu-Mn-Zr alloy, θ′ precipitates are considerably larger at room temperature with higher inter-precipitate distance. Their larger size helps to reduce the coarsening rate on thermal exposure since higher inter-precipitate distance ensures non-overlapping diffusion fields [48].

Furthermore, the constant $k$ in Eq. (4) can be represented as $k=D_{\mathit{Cu}}{\gamma }_{\mathit{sc}}{X}_e$, where $D_{\mathit{Cu}}$ is the diffusional coefficient of Cu in Al, ${\gamma }_{\mathit{sc}}$ is the interfacial energy of semi-coherent interface of θ′ precipitate and $X_e$ is the equilibrium solubility of Cu in Al [52]. Reduction in the interfacial energy of semi-coherent interfaces due to solute segregation therefore helps in decreasing the value of $k$. This in turn contributes to the reduction in the coarsening rate for θ′ precipitates.

As it seems, the microstructural requirement for better coarsening resistance and high temperature stability of Al-Cu-Mn-Zr alloy is quite counterintuitive. At room temperature, a fine-scale microstructure with smaller precipitates and correspondingly, smaller inter precipitate spacing is preferred for high strength [8]. However, a larger precipitate with higher inter-precipitate spacing is desired for enhanced coarsening resistance at higher temperature. Together, Al-Cu-Mn-Zr alloys present low to moderate strength at room temperature but excellent retention of that strength at elevated temperature [48].

8.4 Role of trace elements (Si and Ti)

Other than the major micro-alloying elements (Mn & Zr), trace elements (e.g. Si and Ti) present in the composition may further influence the microstructural stability of Al-Cu-Mn-Zr alloy at elevated temperature. Silicon decreases the coarsening resistance for ${\mathrm{\theta }}^{\prime }$ precipitates so that Si content in Al-Cu-Mn-Zr alloys should be preferably below 0.1 wt% [48]. Si being a faster diffusing species than Zr and even than Mn, preferentially occupies atomic positions at θ′ precipitate/α-Al matrix interfaces above this critical concentration (>0.1 wt%), thereby preventing further segregation of Mn or Zr atoms at these locations. However, Si is not as efficient as Mn or Zr for stabilization of θ′ precipitates at elevated temperatures. Below the threshold concentration, presence of Si is not detrimental for Al-Cu-Mn-Zr alloys as shown by Shower et al. [63]. Si content in the range of 0.05 wt% to <0.1 wt% can even outperform the hardness of base Al-Cu-Mn-Zr alloys at elevated temperature.

After solution treatment of Al-Cu alloys, quench-in vacancies can cluster together to form edge dislocations at room temperature [12, 69]. When these vacancies are in significant density, they can even form dislocation loops rather than individual dislocations, which can further climb and form dislocation helices [70, 71, 72, 73, 74]. These helices accommodate far more number of vacancies with their spacings being larger than individual dislocations. When Si atoms are present in significant quantity in the binary Al-Cu alloy, they can also cluster together during aging due to high diffusivity. The dislocations, dislocation helices and Si clusters can all potentially provide heterogeneous nucleation sites for θ′ precipitates when aged above θ″ solvus.

When the Si content is kept low (<0.05 wt%), Al-Cu alloys essentially act as a binary system and θ′ precipitates mostly nucleate at the dislocation loops, thereby promoting a fine scale microstructure on aging [63]. At higher Si content (0.11 wt%-0.24 wt%), θ′ precipitates nucleate at Si clusters, which again leads to a finer microstructure. However, at the intermediate Si content (0.05–0.1 wt%), nucleation of θ′ precipitates primarily occurs at the dislocation helices. As a result, the number density of θ′ precipitates decreases such that their inter-precipitate spacings become larger as well as the critical size for $\mathrm{\theta }\prime \to \mathrm{\theta }$ transformation increases compared to either high or low Si containing Al-Cu alloys. At room temperature, such coarse microstructure of Si containing Al-Cu alloys yields low hardness in peak-aged condition. However, initially larger θ′ precipitates tend to coarsen far less during elevated temperature exposure since their larger size and greater inter-particle spacings provide better resistance.

Titanium when present in trace concentration can also influence the high temperature stability of Al-Cu-Mn-Zr alloys by forming stable Al₃Ti precipitates having L1₂ crystal structure [75]. Titanium atoms show similar segregation profile as Zr at θ′ precipitate/α-Al matrix interfaces. Poplawsky et al. [62] in this regard observed unique L1₂ structured Al₃(Zr_xTi_1-x) precipitates on the α-Al matrix/ θ′ precipitate interfaces from addition of Ti to Al-Cu-Mn-Zr alloys. The (001) interfaces of this Al₃(Zr_xTi_1-x) precipitates are coherent with ${\left(100\right)}_{\mathit{Al}}$ as well as ${\left(001\right)}_{{\mathrm{\theta }}^{\prime }}$ planes, which helps to stabilize θ′ precipitates by reducing their misfit strain. Such L1₂ precipitate formation further restricts addition of Cu atoms to θ′ precipitates, thereby retarding their thickening along the coherent interfaces. This mechanism of coarsening resistance is known as “ledge poisoning” [63]. The semi-coherent interfaces, on the other hand, are depleted of Al₃(Zr_xTi_1-x) precipitates. This is further confirmed and explained from DFT calculations considering interfacial energetics of preferential precipitation [62, 63]. The formation of Al₃(Zr_xTi_1-x) precipitates on the coherent interface in turn creates θ′/Al₃(Zr_xTi_1-x) /α-Al structure which is energetically favorable compared to similar structures on the semi-coherent interfaces.

Furthermore, self-diffusion coefficients of Mn, Zr and Ti in Al vary in the order D_Mn > D_Zr > D_Ti so that Mn atoms diffuse faster and segregate at the coherent and semi-coherent interface of θ′ precipitates on thermal exposure while slower diffusing Zr atoms segregates primarily at the coherent interfaces [63]. In the peak-aged condition, semi-coherent interface of ${\mathrm{\theta }}^{\prime }$ precipitate is therefore populated only with faster diffusing Mn atoms. On thermal exposure at 300°C for 200 hours, Mn atoms initially segregate at both coherent and semi-coherent interfaces. Afterwards, the slower diffusing Zr atoms segregate at the coherent interfaces. During further heat treatment at 350°C for 200 hours, Ti atoms can diffuse and form Al₃(Zr_xTi_1-x) precipitates on the coherent interfaces as well as at the edges of coherent/semi-coherent interfaces; the relative proportion of these precipitates remains higher on the former location. During the high temperature heat treatment, Mn on the other hand, tend to penetrate towards the bulk of θ′ precipitates. This sequence of segregation of various micro-alloying elements is schematically depicted in Figure 5.

Figure 5.

Schematics showing (a) segregation of Mn at the semi-coherent interface of ${\mathrm{\theta }}^{\prime }$ precipitate in the peak aged condition, (b) Zr segregation after prolonged thermal exposure at 300°C for 200 hours and (c) formation of Al₃(Zr_xTi_1-x) precipitates on the edges of coherent/semi-coherent interfaces as well as penetration of Mn through the bulk of θ′ precipitate after heat treatment at 350°C for 200 hours.

8.5 Computational studies

As discussed before, reduction in interfacial energy due to the segregation of solute atoms (Mn and Zr) at the mobile interfaces promotes thermal stabilization for metastable θ′ precipitates up to a prolonged duration [48]. In this regard, density functional theory (DFT) simulations were carried out to determine the interfacial energy for various θ′ precipitate/α-Al matrix interfaces with and without solute addition (Mn & Zr). In DFT calculations, segregation energy (ΔE_seg) was defined as

$$\Delta E_{\mathit{seg}}=\Delta E_{\mathit{sol}}\left(\mathit{int}\right)-\Delta E_{\mathit{sol}}\left(\mathit{\text{bulk}}\right)$$E5

where, $\Delta E_{\mathit{sol}}\left(\mathit{int}\right)$ and $\Delta E_{\mathit{sol}}\left(\mathit{\text{bulk}}\right)$ is the heat of solution when certain solute element situates at θ′ precipitate/α-Al matrix interface and in the bulk of the precipitate, respectively. The interfacial energy change is then defined as:

$$\mathit{\Delta \gamma }=\frac{\Delta E_{\mathit{seg}}}{2A}$$E6

where, A is the area of the interface. The calculations suggest that the interfacial energy for pure (i.e. without any segregation) coherent and ${\left(100\right)}_{{\mathrm{\theta }}^{\prime }}$ semi-coherent interfaces are 252 mJ/m² and 527 mJ/m², respectively [48]. It further establishes that the addition of Mn and Zr in Al-Cu-Mn-Zr alloys decreases the interfacial energy for both coherent and semi-coherent interfaces. It also corroborates with the experimental observations that Mn atoms preferentially segregate at the semi-coherent interfaces while Zr atoms occupies places both at coherent and semi-coherent interfaces [62].

Important to note that the coarsening of strengthening precipitates is NOT a function of interfacial energy alone, although reduction in interfacial energy certainly adds as a dominating factor for hindering the precipitate coarsening [76, 77]. Lattice misfit strain is another important driving force for coarsening of θ′ precipitates. Other kinetic factors like diffusion barrier formation and solute drag are crucial too for restricting coarsening of θ′ precipitates on thermal exposure [52, 78]. Due to the formation of diffusion barrier around θ′ precipitates, mobility of Cu atoms within α-Al matrix reduces to a large extent, which further helps to prevent coarsening or transformation of these precipitates. In addition, the presence of slow diffusing element/s in any Al alloy system can contribute to the coarsening resistance of θ′ precipitate [79]. Such slower diffusing element creates a solute drag since they need to be displaced during the growth of the interfaces, which in turn retards the coarsening of θ′ precipitate.

Introduction of a third element (Mn, Zr etc.) within the binary Al-Cu alloys can lead to one or a combination of thermodynamic and kinetic restrictions (mentioned above) to precipitate coarsening processes. Shower et al. [76] in this regard carried out a phase field modeling study to understand the synergistic effect of various mechanisms that offers precipitate coarsening resistance. The study suggests that a combination of interfacial energy reduction and solute drag due to the addition of Mn and Zr contributes to the coarsening resistance of θ′ precipitates up to 300°C. In the process, a continuous segregation profile forms for Mn atoms along the interfaces of θ′ precipitates with a larger weightage at the semi-coherent interface. Other solute atoms e.g. Zr, which introduce a positive misfit strain do not effectively interact with the mobile semi-coherent interfaces; rather they tend to segregate at the coherent interfaces. When the working temperature is raised to 400°C, resistance to precipitate coarsening requires simultaneous reduction in the mobility of Cu atoms through α-Al matrix and the interfacial energy of θ′ precipitates. This could not be achieved by micro-alloying with Mn and Zr alone, which is why the Al-Cu-Mn-Zr alloys losse their excellent thermal stability at 400°C and beyond. All the cumulative and inter-connected effects that contribute to the stabilization of θ' precipitate at elevated temperature in Al-Cu-Mn-Zr alloy are schematically shown in Figure 6.

Figure 6.

Schematic flowchart showing the cumulative effect of various contributing factors and mechanisms involved in the stabilization of θ′ precipitates at elevated temperatures for Al-Cu-Mn-Zr alloys.

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′ 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′ 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.

Figure 7.

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

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 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 inter-precipitate 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, 82, 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.

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′ 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.

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].

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 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′ precipitates. In this temperature range, 319 alloy losses its flow stress by 40%. At lower strain rates (e.g. 10⁻⁴ s⁻¹ & 10⁻³ 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′ 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).

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⁻¹ strain rate) SEM micrograph of RR350 alloy showing shear band formation and bending of θ′ precipitates within α-Al grains having <100> direction nearly parallel to the compression axis [87].

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 10⁵ 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 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′ 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.

11. Industrial application potentials of Al-Cu-Mn-Zr alloys

Industrial application of Al-Cu-Mn-Zr alloys require certain additional considerations on and above their excellent thermal stability and associated improvement in most of the high temperature mechanical properties as mentioned above. These include, but certainly not restricted to the ease of casting and defect formation, assessment of mechanical properties e.g. fatigue testing at larger component scale, possibility to adopt alternate component fabrication methodologies like additive manufacturing, wrought processing etc. [89]. Some of these aspects are mentioned below:

11.1 Hot tearing resistance

Hot tearing is a crucial casting defect that can affect the structural stability and properties of as-cast components [90, 91]. During solidification, molten metal usually remains in semi-solid state (mushy zone) for considerable duration. It also undergoes severe volume contraction and associated thermal stresses within the semi-solid metal regions. Under this condition, cracks form in the solidified component if there is an inadequate supply of molten mass to fill up the shrinkage volume. Controlling such defects in castings is difficult but extremely important for improving the fatigue life per se [91].

Sabau et al. [92] studied the hot tearing resistance of cast Al-Cu-Mn-Zr alloys with and without grain refiners in comparison to base Al-Cu and Al-Si alloys. The base Al-Cu alloy with >7 wt% Cu exhibits a grain refined microstructure in the casting. It also shows a decrease in the length of columnar to equiaxed transition zone that in turn improves the hot tearing resistance. In addition, simultaneous presence of Si and Fe (>0.2 wt%) increases the hot tearing resistance for this alloy. The low Cu containing alloys, on the other hand, possess coarse columnar grains within the cast microstructures which facilitates hot tearing for them. For the Al-Cu-Mn-Zr alloys, when Cu is added above 7 wt%, significant grain refinement occurs which further contributes to their excellent hot tearing resistance. In addition, when 0.1 wt% Ti is added as additional grain refiner, resultant Al-Cu-Mn-Zr alloy exhibits the finest microstructure and correspondingly, the best hot tearing resistance [89]. Ti added Al-Cu-Mn-Zr alloy was therefore speculated suitable for industrial applications [92].

In the measurement of hot tearing resistance, Sabau et al. [92] used an in-house multi-arm casting setup with varying arm length in a permanent mold. In this six-armed mold, the shortest arms were free from any visible cracks for all the alloys having varying amount of Cu and Ti, while the longest arm had severe cracking for these alloys. As per the visible inspection, a cracking index (C_i) was assigned to each of the arms where 0 corresponds to no crack condition and 2.5, 5, 7.5, and 10 was assigned to small, moderate, severe and completely fractured arms, respectively. Hot tearing index (HTI), M_s was then defined as the simple average of C_i over these six arms i.e.

$$M_s=\frac{1}{6}\sum _{i=1}^6{C}_i$$E7

The length of the arms, however, also plays a crucial role in the formation of these cracks. The longer arms are more susceptible to cracking compared to the shorter ones so that the weighted average was preferred for the calculation of HTI. The weighted average (M_J-K) was defined as:

$$M_{J-K}=\frac{{\sum }_{i=J}^K{w}_i{C}_i}{{\sum }_{i=J}^K{w}_i}$$E8

where, J and K corresponds to the shortest and longest arms and w_i is the weight factor which is inversely proportional to the length of the arm; J was considered as 2 and K as 5. As per the industrial standard, hot tearing index (HTI) equal or less than 3.3 is considered acceptable when C_i is in the range of 0–10. The HTI for various Al-Cu-Mn-Zr alloys with promising industrial application potential are listed in Table 1. Important to note that the fluidity of Al-Cu-Mn-Zr alloys with respect to temperature, especially under solidification conditions is also of immense importance for successful industrial application standpoint. However, no such data reporting the fluidity of Al-Cu-Mn-Zr alloys is available yet.

Alloy	Ti content (wt%)	$M_{J-K}$
Al-Cu-Mn-Zr with 8 wt% Cu	0.02	2.6
	0.1	2.01
	0.2	2.36
Al-Cu-Mn-Zr with 6.2 wt% Cu	0.1	2.6
	0.2	3.21
Al-Cu-Mn-Zr with 7.3 wt% Cu	0.1	2.49
	0.2	2.84

Table 1.

Hot tearing index (HTI) of different Al-Cu-Mn-Zr alloys with varying Ti content.

11.2 Additive manufacturing of Al-Cu-Mn-Zr alloy

In recent times, additive manufacturing (AM) is proven to be an extremely useful and alternate technique for shaping intricate parts in industrially relevant scales with excellent property combinations compared to cast counterparts [93]. For Al alloys however, additive manufacturing is a rather complicated and challenging process due to several factors like poor powder flowability, high thermal conductivity, laser reflectivity etc. [94]. In this regard, Shyam et al. [95] successfully fabricated AM parts from Al-Cu-Mn-Zr alloys by selective laser melting (SLM) without any hot tearing using optimized processing parameters. The substrate temperature for SLM was kept 200°C which was sufficient for in-situ formation of strengthening θ′ precipitates.

Due to laser melting, AM Al-Cu-Mn-Zr alloys form typical “peacock tail” microstructure having overlapping melt pools. In addition, the AM microstructure consists of long columnar grains at the top and fine equiaxed grains at the bottom of the melt pools. Such refined AM microstructure yields comparatively higher strength up to 300°C compared to the cast Al-Cu-Mn-Zr alloys. The bimodal grain size distribution and refined grain boundary intermetallic precipitates further enhance the tensile elongation for AM Al-Cu-Mn-Zr alloys. However, creep properties for these AM alloys are somewhat compromised compared to the cast counterparts due to high proportion of grain boundaries in the refined AM microstructure. Overall, AM Al-Cu-Mn-Zr alloys are envisioned having potential in complex component manufacturing for high temperature applications owing to the simultaneous positive effects of refined microstructure and in situ formation of thermally stable strengthening θ′ precipitates [95].

11.3 Environmental impact from Al-Cu-Mn-Zr alloy

In the current global scenario, any new alloy development must help in reducing environmental impact e.g. carbon footprint and green house emissions [96]. The primary target area for Al-Cu-Mn-Zr alloys is automotive industry, which also formed the early motivation of their inception and further development [97]. The aim was to develop Al alloys for engine components that experience high working temperature (∼300°C) e.g. cylinder heads in combustion engines. The use of Al-Cu-Mn-Zr alloys for making such components can effectively raise the working temperature and increase the fuel efficiency of next generation passenger vehicles, thereby proving environment friendly in terms of fuel consumption.

12. Conclusion

The present book chapter elucidates a comprehensive review about the development as well as the science and technology behind the new-age Al-Cu-Mn-Zr (ACMZ) alloys. The major observations are summarized below.

• The new age Al-Cu-Mn-Zr (ACMZ) alloy developed in recent times by micro-alloying conventional Al-Cu alloys with Mn and Zr shows unprecedented microstructural stability up to ∼350°C.

• The strengthening θ′ precipitates in Al-Cu-Mn-Zr alloys are stabilized primarily due to preferential solute segregation (Mn & Zr) at the θ′ precipitate/α-Al matrix interfaces which helps in reducing their interfacial energy. Mn atoms segregate at the mobile semi-coherent interfaces while Zr atoms primarily sits at the corner of coherent/semi-coherent interfaces.

• The difference in the diffusivity of micro-alloying elements decides the sequence of their segregation at θ′ precipitate/α-Al matrix interfaces. This in turn plays a synergistic role in stabilizing θ′ precipitate at different temperature ranges. Faster diffusing Mn initially segregates at both coherent and semi-coherent interfaces and provides thermal stability at lower temperatures. Zr being relatively slower diffusing element segregates at a later stage but effectively restricts ${\mathrm{\theta }}^{\prime }$ precipitates from transforming to stable θ precipitates up to a much higher temperature.

• Various other mechanisms e.g. solute drag, diffusion barrier formation, ledge poisoning by co-precipitation of thermally stable intermetallics etc. also contributes to the coarsening resistance of ${\mathrm{\theta }}^{\prime }$ precipitates.

• Trace elements like Ti and Si further influence θ′ precipitate stability in Al-Cu-Mn-Zr alloys. Si above a critical content (>0.1 wt%) potentially substitutes Mn and Zr at the θ′ precipitate/α-Al matrix interfaces, thereby decreasing the stability of ${\mathrm{\theta }}^{\prime }$ precipitates at elevated temperatures. On the other hand, L1₂ structured Al₃(Zr_xTi_1-x) precipitates form at θ′ precipitate interfaces in presence of Ti which further adds to the stability of these alloys at higher temperatures.

• DFT simulations confirm the reduction of interfacial energy from segregation of micro-alloying elements and provide the segregation profiles at various (coherent and semi-coherent) interfaces. Phase field simulations further suggest that the cumulative effect of interfacial energy reduction and solute drag led to the θ′ precipitate stabilization at 300°C. At 400°C, reduction in the mobility of Cu atoms is required in addition to reduced interfacial energy for efficient resistance towards θ′ precipitate coarsening.

• As a result of the excellent thermal stability of strengthening ${\mathrm{\theta }}^{\prime }$ precipitates, Al-Cu-Mn-Zr alloys exhibit superior high temperature tensile properties (hardness, yield and ultimate tensile strengths etc.) compared to base Al-Cu or Al-Si-Cu alloys. For example, Al-Cu-Mn-Zr alloys possess almost twice the yield strength compared to base Al-Cu alloy after prolonged thermal exposure at 300°C.

• At high temperature (300–350°C), compressive flow stress of Al-Cu-Mn-Zr alloy (RR350) is higher than conventional Al-Si-Cu (319) alloy. The high temperature deformation mechanisms for the former alloy vary from strain hardening at room temperature to cross slip and dynamic recovery as the test temperature increases at 250°C. Further higher test temperature promotes dynamic recrystallization of the matrix at 300°C and finally lead to grain boundary sliding to 350°C.

• The Al-Cu-Mn-Zr alloy shows excellent creep resistance compared to the base Al-Cu alloys. For the former alloy, diffusional Coble creep is the controlling mechanism at low stress levels which switches to dislocation creep at higher stress levels.

• In high temperature low cycle fatigue testing, Al-Cu-Mn-Zr alloy does not fail up to 10⁵ cycles at low stress amplitudes whereas the fatigue life decreases with increasing stress amplitude. The casting defects (pores) are found to be more influential factor by means of stress concentration and crack initiation than either the stability of ${\mathrm{\theta }}^{\prime }$ precipitates or the presence of grain boundary intermetallic precipitates.

• The Al-Cu-Mn-Zr alloy exhibits superior hot tearing resistance compared to the conventional Al-Cu alloys. Additive manufacturing of Al-Cu-Mn-Zr alloy also shows promising property combinations owing to a refined microstructure and in-situ formation of strengthening ${\mathrm{\theta }}^{\prime }$ precipitates. In addition to the positive environmental impacts, both of these factors are crucial for their suitability in industrial production purpose.

Overall, Al-Cu-Mn-Zr alloys offer huge potential for industrial applications where lightweight materials are sought after for excellent high temperature mechanical properties. These new-age alloys can in fact prove to be a game changer for the existing passenger vehicle engines and may bring a paradigm shift in the automobile sectors. Owing to low density and excellent high temperature stability, their future use is certainly leading to a safer and greener environment.