**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<sup>0</sup> 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<sup>0</sup> , 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<sup>0</sup> precipitates coarsen due to the enhanced diffusion of solute Cu atoms [8]. Segregation of Mn and Zr atoms at the θ<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> precipitates retain their morphology and aspect ratio on high temperature heat treatment in case of Al-Cu-Mn-Zr alloys.

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<sup>0</sup> 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.

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

#### **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].*

The stability of θ<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> 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].

In order to understand the individual and synergistic effect of Mn and Zr microalloying 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

#### **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<sup>0</sup> *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].*

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<sup>0</sup> 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 θ<sup>0</sup> precipitates unless interfacial segregation resists [63]. Furthermore, self-diffusion coefficient of Zr in

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

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 θ<sup>0</sup> 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 θ<sup>0</sup> 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), θ<sup>0</sup> 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/θ<sup>0</sup> precipitates. However, sufficient coarsening of θ<sup>0</sup> 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<sup>0</sup> 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 θ<sup>0</sup> 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<sup>0</sup> 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]:

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

where, *r* & *r*<sup>0</sup> are the mean precipitate radius at time *t* ¼ *t* and *t* ¼ 0, respectively and *k* is a constant. Partial derivation of Eq. (3) with respect to *t* gives*:*

$$\frac{\partial \overline{r}}{\partial t} = \frac{k}{\Im \left( k.t + \overline{r}\_0^{\;\;\;\;\phi} \right)^{\frac{2}{5}}} \tag{4}$$

Since *<sup>∂</sup><sup>r</sup> <sup>∂</sup><sup>t</sup>* varies with the initial precipitate radius as � <sup>1</sup> *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<sup>0</sup> 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* ¼ *DCuγscXe*, where *DCu* is the diffusional coefficient of Cu in Al, *γsc* is the interfacial energy of semi-coherent interface of θ<sup>0</sup> precipitate and *Xe* 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<sup>0</sup> 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 θ<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> 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 [70–74]. 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<sup>0</sup> precipitates when aged above <sup>θ</sup>″ solvus.

When the Si content is kept low (<0.05 wt%), Al-Cu alloys essentially act as a binary system and θ<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> precipitates primarily occurs at the dislocation helices. As a result, the number density of θ<sup>0</sup> precipitates decreases such that their inter-precipitate spacings become larger as well as the critical size for θ<sup>0</sup> ! θ 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 peakaged condition. However, initially larger θ<sup>0</sup> 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

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

θ<sup>0</sup> precipitate/α-Al matrix interfaces. Poplawsky et al. [62] in this regard observed unique L12 structured Al3(ZrxTi1-x) precipitates on the α-Al matrix/ θ<sup>0</sup> precipitate interfaces from addition of Ti to Al-Cu-Mn-Zr alloys. The (001) interfaces of this Al3(ZrxTi1-x) precipitates are coherent with 100 ð Þ*Al* as well as 001 ð Þ<sup>θ</sup><sup>0</sup> planes, which helps to stabilize θ<sup>0</sup> precipitates by reducing their misfit strain. Such L12 precipitate formation further restricts addition of Cu atoms to θ<sup>0</sup> 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(ZrxTi1-x) precipitates. This is further confirmed and explained from DFT calculations considering interfacial energetics of preferential precipitation [62, 63]. The formation of Al3(ZrxTi1-x) precipitates on the coherent interface in turn creates θ<sup>0</sup> /Al3(ZrxTi1-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 *DMn > DZr > DTi* so that Mn atoms diffuse faster and segregate at the coherent and semi-coherent interface of θ<sup>0</sup> 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 θ<sup>0</sup> 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(ZrxTi1-x) precipitates on the coherent interfaces as well as at the edges of coherent/semicoherent 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<sup>0</sup> precipitates. This sequence of segregation of various micro-alloying elements is schematically depicted in **Figure 5**.

#### **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<sup>0</sup> precipitates up to a prolonged duration [48]. In this regard, density functional theory (DFT) simulations were carried out to determine the

#### **Figure 5.**

*Schematics showing (a) segregation of Mn at the semi-coherent interface of* θ<sup>0</sup> *precipitate in the peak aged condition, (b) Zr segregation after prolonged thermal exposure at 300°C for 200 hours and (c) formation of Al3(ZrxTi1-x) precipitates on the edges of coherent/semi-coherent interfaces as well as penetration of Mn through the bulk of* θ<sup>0</sup> *precipitate after heat treatment at 350°C for 200 hours.*

interfacial energy for various θ<sup>0</sup> precipitate/α-Al matrix interfaces with and without solute addition (Mn & Zr). In DFT calculations, segregation energy *(ΔEseg)* was defined as

$$
\Delta E\_{\text{seg}} = \Delta E\_{\text{sol}}(\text{int}) - \Delta E\_{\text{sol}}(\text{bulk}) \tag{5}
$$

where, *ΔEsol*ð Þ *int* and *ΔEsol*ð Þ *bulk* is the heat of solution when certain solute element situates at θ<sup>0</sup> precipitate/α-Al matrix interface and in the bulk of the precipitate, respectively. The interfacial energy change is then defined as:

$$
\Delta \chi = \frac{\Delta E\_{\text{seg}}}{\Delta A} \tag{6}
$$

where, *A* is the area of the interface. The calculations suggest that the interfacial energy for pure (i.e. without any segregation) coherent and 100 ð Þ<sup>θ</sup><sup>0</sup> semi-coherent interfaces are 252 mJ/m<sup>2</sup> and 527 mJ/m<sup>2</sup> , 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<sup>0</sup> precipitates. Other kinetic factors like diffusion barrier formation and solute drag are crucial too for restricting coarsening of θ<sup>0</sup> precipitates on thermal exposure [52, 78]. Due to the formation of diffusion barrier around θ<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> 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<sup>0</sup> precipitates up to 300°C. In the process, a continuous segregation profile forms for Mn atoms along the interfaces of θ<sup>0</sup> 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<sup>0</sup> 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**.

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

**Figure 6.**

*Schematic flowchart showing the cumulative effect of various contributing factors and mechanisms involved in the stabilization of* θ<sup>0</sup> *precipitates at elevated temperatures for Al-Cu-Mn-Zr alloys.*
