**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<sup>0</sup> *)* undergo rapid growth and coarsening and even transform to stable incoherent precipitates (metastable θ<sup>0</sup> ! 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 microalloying with Mn and Zr which possesses excellent stability for strengthening θ<sup>0</sup> 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 GP -I zones (Guinier Preston zone) along a plane parallel to the (001) plane of the α-Al matrix.

When the aging process continues above the GP zone solvus line (**Figure 1a**), θ<sup>00</sup> metastable precipitates (also known as GP-II zones) form by diffusion of Cu atoms from the solute cluster zones (GP -I zone). θ<sup>00</sup> are characterized by the chemical formula Al3Cu and ordered tetragonal structure (*a* ¼ *b* ¼ 4*:*04Å,*c* ¼ 7*:*68Å) having two layers of Cu atoms separated by three layers of Al atoms (**Figure 1b**) [10]. It is completely

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

**Figure 1.**

*Schematics showing (a) section of Al-Cu binary phase diagram, (b) crystal structure of parent α-Al matrix and various metastable (*θ<sup>00</sup> *and* θ<sup>0</sup> *) and stable (*θÞ *precipitates in Al-Cu alloy system, (c) morphology and different interfaces for metastable* θ<sup>0</sup> *precipitate.*

coherent with α-Al matrix from all sides although misfit strain develops along ð Þ 100 *Al* and 010 ð Þ*Al* planes. Metastable θ<sup>0</sup> precipitates next form from θ<sup>00</sup> precipitates on further aging. θ<sup>0</sup> has a stoichiometric formula of Al2Cu with body centred tetragonal structure (*a* ¼ *b* ¼ 4*:*04Å,*c* ¼ 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 θ00!θ<sup>0</sup> transformation, the later precipitate starts nucleating at the preexisting θ<sup>00</sup> precipitates via continued diffusion of Cu atom and continues to grow until the entire θ<sup>00</sup> precipitate transforms into θ<sup>0</sup> . In this transformation, θ<sup>0</sup> holds the same orientation relationship with α-Al matrix as does θ<sup>00</sup> precipitate i.e. ð Þ 001 <sup>θ</sup>0k ð Þ 001 *Al* and 100 ½ �θ0k ½ � 100 *Al* [12]. It grows as a plate within the α-Al matrix parallel to the 010 ð Þ*Al* and remains coherent along this plane. However, along the ð Þ 100 *Al* and 010 ð Þ*Al* planes, θ<sup>0</sup> becomes incoherent or complex semi-coherent with α-Al matrix. The semi-coherent side of the plate shaped θ<sup>0</sup> precipitate is not completely circular; rather it forms as octagon with facets along 100 ½ �<sup>θ</sup><sup>0</sup> and 110 ½ �<sup>θ</sup><sup>0</sup> directions (**Figure 1c**). The 110 f g<sup>θ</sup><sup>0</sup> interface edges further act as a solute gateway from where Cu atoms diffuse and coarsen θ<sup>0</sup> precipitates on prolonged thermal exposure. The atomic arrangement as well as ledge dislocations on 110 f gθ<sup>0</sup> semicoherent interfaces further assist in the accommodation of Cu atoms [13].

The equilibrium precipitate in the Al-Cu system is tetragonal θ (*a* ¼ *b* ¼ 6*:*07Å,*c* ¼ 4*:*87Å, space group I4/mcm) having chemical formula Al2Cu (**Figure 1b**) [10]. It is incoherent with the α-Al matrix along all sides. θ forms from θ<sup>0</sup> precipitate via mechanism similar to θ<sup>00</sup> ! θ<sup>0</sup> transformation. The habit plane of θ 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: Solute clusters ! GP zones ! GP II zones θ<sup>00</sup> ð Þ! θ<sup>0</sup> ! θ*:*

## **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. GP -I or θ00), 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 θ<sup>0</sup> ) are large in size and possess semi-coherent or incoherent interfaces with α-Al matrix, the dislocation line bulges within the interprecipitate 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:

$$
\Delta \tau = \left(\frac{3}{2\pi}\right)^{\frac{1}{2}} \frac{Gb}{r} \cdot f^{\frac{1}{2}} \tag{1}
$$

where, Δ*τ* 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 θ<sup>0</sup> precipitate forming on {100}Al habit plane having diameter *D* and thickness *t*, CRSS increment, Δ*τ* due to Orowan looping is given by [17]:

$$
\Delta \tau = \left(\frac{Gb}{2\pi\sqrt{1-\nu}}\right) \left(\frac{1}{0.931\sqrt{\frac{0.306\pi D T}{f} - \frac{\pi D}{t} - 1.061T}}\right) \left(\ln\frac{1.225t}{r\_0}\right) \tag{2}
$$

where, *ν* is the Poisson's ratio of α-Al matrix and *r*<sup>0</sup> 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 θ<sup>00</sup> 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 θ<sup>00</sup> ! θ<sup>0</sup> transformation. Afterwards, dislocation bowing around large and semi-coherent θ<sup>0</sup> precipitates is the main strengthening process. On further coarsening, the number density continuously decreases for θ<sup>0</sup> precipitates so that the alloy

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

#### **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].*

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 GP -I and/or θ<sup>00</sup> 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 GP-I and θ<sup>00</sup> precipitates is the primary strengthening mechanism here. As the aging time increases, GP -I zones continuously transform to θ<sup>00</sup>, thereby increasing their number density in the α-Al matrix. Some of them even transform to other metastable precipitates e.g. θ<sup>0</sup> . 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 θ<sup>0</sup> 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 θ<sup>0</sup> precipitate size, aspect ratio, inter-precipitate distance etc. are now optimum for maximum hindrance towards dislocation motion due to shearing through the precipitates.

The θ<sup>0</sup> precipitates continuously coarsen with increasing aging time and even start transforming to stable θ precipitate under prolonged aging [12]. The energy required to shear θ<sup>0</sup> 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 θ<sup>0</sup> 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 overaging, 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.
