Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys

*Jacobo Hernandez-Sandoval, Mohamed H. Abdelaziz, Agnes M. Samuel, Herbert W. Doty and Fawzy H. Samuel*

## **Abstract**

The present study focused on the tensile properties at ambient and high temperatures of alloy 354 without and with the addition of zirconium. Tensile tests were performed on alloy samples submitted to various aging treatments, with the aim of understanding the effects of the addition made on the tensile properties of the alloy. Zirconium reacts only with Ti, Si, and Al in the alloys examined to form the phases (Al,Si)2(Zr,Ti) and (Al,Si)3(Zr,Ti). Testing at 25°C reveals that the minimum and maximum quality index values, 259 and 459 MPa, are observed for the as-cast and solution heat-treated conditions, respectively. The yield strength shows a maximum of 345 MPa and a minimum of 80 MPa within the whole range of aging treatments applied. The ultimate tensile and yield strength values obtained at room temperature for T5-treated samples stabilized at 250°C for 200 h are comparable to those of T6-treated samples stabilized under the same conditions, and higher in the case of elevated-temperature (250°C) tensile testing. Coarsening of the strengthening precipitates following such prolonged exposure at 250°C led to noticeable reduction in the strength values, particularly the yield strength, and a remarkable increase in the ductility values.

**Keywords:** aluminum alloys, aging, thermal exposure, tensile testing, precipitation, fractography

## **1. Introduction**

The 354 alloy belongs to the Al-Si-Cu-Mg system similar to B319 alloy that is widely used for automotive engine blocks [1]. The high silicon content in the 354 alloy improves the alloy castability whereas the presence of Cu and Mg noticeably enhances the yield strength (YS) and the ultimate tensile strength (UTS) of the 354 alloy due to the formation of intermetallic phases, mainly Al2Cu or eutectic Al + Al2Cu, and Mg2Si precipitates [2, 3]. However, segregation behavior of Cu may lead to incipient melting during solution treatment which will apparently reduce the alloy strength [4]. Addition of Mg has a strong affinity to react with Sr, leading to the formation of a complex Mg2SrAl4Si3 intermetallic phase, and hence reducing the effectiveness of Sr as a Si modifying agent [5]. In the absence of Cu, high Fe and Mg contents lead to the formation of π-FeMg3Si6Al8 phase which is difficult to dissolve during the solution

treatment process [6, 7]. In the quaternary Al-Si-Cu-Mg alloy system, *Q*-phase (Al4Mg8Cu2Si6) can coexist with the Al2Cu, Mg2Si, and Si phases depending on the levels of Cu, Mg, and Si [8–11]. The different factors that may influence the mechanical behavior of cast aluminum alloys are schematically represented in **Figure 1** [12].

Zirconium may be added to Al alloys in order to refine the grain structure due to the presence of fine coherent dispersoids (mainly Al3Zr) which obstruct dislocation motion and in turn, enhance the elevated temperature mechanical properties of aluminum alloys [13]. In order to increase the volume fraction of Al3Zr precipitates and based on the phase diagram of Al-Zr, the concentration of Zr in the alloys investigated in this study was kept at around 0.3 wt.% [14].

The main purpose of solution heat treatment is to obtain a supersaturated solid solution at high temperatures (below the eutectic temperature). As a result, a homogeneous supersaturated solid solution (SSSS) will form through dissolving the precipitated phases during the solidification process, such as β-Mg2Si, θ-Al2Cu, Q-Al5Cu2Mg8Si6, π-Al9FeMg3Si5 and β-Al5FeSi phases. The β-Mg2Si and θ-Al2Cu phases can be easily dissolved when the optimum solution heat treatment temperature and time are employed. The solution treatment temperature is determined according to the alloy composition and solid solubility limit; however, it must be lower than the melting point of the phases that exist in the as-cast structure to avoid incipient melting of these phases [15, 16].

Following the empirically developed concept of quality index proposed by Drouzy et al. [17, 18] Cáceres proposed a mathematical model emphasizing the significance of the quality index as follows [17, 19, 20]:

$$Qc = \left[ \left( qn \right)^{\pi} \exp^{-qn} + \text{0.4} \log \left( \mathbf{100} qn \right) \right] \tag{1}$$

where quality index *Q* can be calculated using the relative quality index (*q*), strain-hardening exponent (*n*), and the strength coefficient (K).

#### **Figure 1.**

*Schematic representation of factors affecting alloy performance [12].*

*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

The present study was undertaken to explore the effect of Zr addition and aging conditions of the as cast tensile bars on:


It should be noted here that the term "temperature" applies to aging temperatures as well as testing temperature.

## **2. Experimental procedure**

Alloy 354 modified with 200 ppm of strontium (using Al-10% Sr master alloy) and grain refined using 0.20 wt.%Ti (Al-5%Ti-1%B) was used as the base alloy (alloy A). To this alloy, 0.3%Zr in the form of Al-25wt.%Zr master alloy was added (alloy B). The chemical compositions of both alloys are listed in **Table 1**. **Figure 2** shows the microstructure of the as-received base alloy ingots. Melting and casting procedures were carried out as described elsewhere.


**Table 1.**

*Chemical composition of the 354 alloys used in this study.*

**Figure 2.** *Microstructure (200×) of the base alloy 354 used in this work.*

To prepare test bars for the tensile tests, three samples for chemical analysis were also taken at the time of the casting; this was done at the beginning, in the middle, and at the end of the casting process to ascertain the exact chemical composition of each alloy. The experimental work was divided into two stages: Stage I in which the 354 alloy (alloy A) was used, and Stage II where the 354 alloy with 0.3%Zr (alloy B) was used. In Stage I, the melt temperature was kept around 750°C, whereas in Stage II, the melt temperature was superheated to 800°C, to ensure the complete decomposition of the Al-25%Zr master alloy used.

## **2.1 Stage I-alloy A**

Tensile bars were solution heat treated at 495°C for 8 h, followed by quenching in warm water at 60°C, after which artificial aging was applied according to the plan listed in **Table 2**. After aging, the test bars were allowed to cool naturally at room temperature (25°C). All of the samples, whether as-cast, solution heat-treated, or aged, were tested to the point of fracture using an MTS servo-hydraulic mechanical testing machine at a strain rate of 4 × 10<sup>−</sup><sup>4</sup> s<sup>−</sup><sup>1</sup> .

The yield strength (YS) was calculated according to the standard 0.2% offset strain, and the fracture elongation was calculated as the percent elongation (%El) over 50 mm gauge length, as recorded by the extensometer. The ultimate tensile strength (UTS) was also obtained from the data acquisition system of the MTS machine. The average %El, YS, or UTS values obtained from the five samples tested per condition were considered to be the values representing that specific condition. An extensometer, or strain gage was used in the tests to measure the extent of deformation in the samples.

Samples for metallography were sectioned from the tensile-tested bars of all the alloys studied, about 10 mm below the fracture surface. The percentage porosity and eutectic Si-particle characteristics were measured and quantified using an optical microscope linked to a Clemex image analysis system. The microstructures of the polished sample surfaces were examined using an Olympus PMG3 optical microscope. Phase identification was carried out using electron probe microanalysis (EPMA) in conjunction with wavelength dispersive spectroscopic (WDS) analysis, using a JEOL\*JXA-889001WD/ED combined microanalyzer operating at 20 kV and 30 nA, where the electron beam size was ~2 μm.

Mapping of certain specific areas of the polished sample surfaces was also carried out where required, so as to show the distribution of different elements within the phases. The fracture surfaces of tensile-tested samples were also examined using the same SEM, employing the backscattered electron (BSE) detector and EDS system. The fracture behavior was analyzed using the backscattered electron (BSE) images


**Table 2.** *Artificial aging conditions used for room temperature tension tests.* *Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

obtained, and analysis of the EDS spectra of phases observed on the fracture surface. Differential scanning calorimetry (DSC) was used to characterize the sequence of reactions occurring during the heating and/or cooling cycles of an alloy sample during a DSC scan which continuously changes with the increasing or decreasing temperature cycle to produce peaks according to the two expected reactions:


## **2.2 Stage II-alloy B**

For the high temperature tensile tests, samples from selected conditions were tested to fracture using an Instron Universal mechanical testing machine at a strain rate of 4 × 10<sup>−</sup><sup>4</sup> s<sup>−</sup><sup>1</sup> . The heating furnace installed on the testing machine is an electrical resistance, forced-air box type, having the dimensions 30 × 43 × 30 cm. The yield strength (YS) was calculated according to the standard 0.2% offset strain, and the fracture elongation was calculated as the percent elongation (%El) over the 25.4 mm gauge length as recorded by the extensometer. The ultimate tensile strength (UTS) was obtained from the data acquisition system of the universal machine. In order to reach and stabilize the intended test temperature during the tests, at the time that the samples were mounted in the tensile machine, the furnace was already pre-set at the required temperature; also, these samples were kept mounted in the furnace of the tensile testing machine for 30 min before the start of every test.

## **3. Results and discussion**

#### **3.1 Stage I-alloy A**

**Figure 3** shows the macrostructure revealing the grain size for alloy A, about 200 μm. A complete modification of the silicon particles in the microstructure of alloy A in the as-cast condition can be seen in **Figure 4(a)**. From **Figure 4(a)** and **(b)**, solution heat treatment has changed the morphology of the silicon particles from faceted to globular. As a consequence of solution heat treatment, there may also be observed a reduction in the number of silicon particles and a reduction in the density of the silicon phase, due to the diffusion of silicon into the aluminum matrix. The white arrows in **Figure 4(a)** show the rounded shape of the dendrites with grain refining [21], whereas **Figure 4(b)** reveals the dissolution of the Al2Cu phase observed in **Figure 4(a)**—circled.

Zhu and Liu [22] proposed a model of the granulation of unmodified eutectic Si composed of three major stages during heat treatment: (i) the mass transport of solute, (ii) a discontinuous phase fragmentation, and lastly (iii) spheroidization. During heat treatment, the silicon atoms in the matrix at the Si particle tips diffuse to locations on the curved surfaces of the particles, leading to the dissolution of eutectic silicon at the tips. This transport of silicon atoms ultimately causes the fragmentation and spheroidization of eutectic silicon which is important from strength point of view compared to Si particles with sharp edges which act as sites for stress concentration.

The values of secondary dendrite arm spacing (SDAS), porosity, modification level, and grain size for both the as-cast (AC) and solution heat-treated (SHT) condition are listed in **Tables 3** and **4**. As can be seen, SHT resulted in (i) no noticeable change in both the SDAS and grain size, (ii) a significant decrease in the particle density due to coarsening of the eutectic Si particles, and (iii) almost

complete solubility of Al2Cu in the aluminum matrix. Since the solutionizing temperature was well below the incipient melting temperature, tensile test bars revealed negligible change in the amount of porosity, i.e., no incipient melting.

#### **Figure 3.**

*Macrograph showing grain size of the tensile bars in the as-cast condition.*

**Figure 4.** *Optical microstructure: (a) before, and (b) after solution heat treatment.*


#### **Table 3.**

*SDAS, porosity%, grain size, level of modification, and volume fraction of intermetallics for alloy A.*


*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

**Table 4.**

*Silicon particle characterization for alloy A.*

**Figure 5.**

*Variation in alloy tensile parameters as a function of aging temperature and time: (a) UTS, (b) YS, and (c) %El.*

**Figure 5** illustrates the effect of aging treatment on the alloy strength parameters. The main observations inferred from this figure can be summarized as follows:


In order to analyze the alloy quality by means of the Quality Index charts**,** the ascast and the solution heat treated conditions plus aging conditions at 155°C, 190°C, and 350°C for aging times in the range of 2–100 h were used. From a previous study [5], K was calculated as 500 MPa.

The plastic strain and the quality index (*Q*) both exhibit a great improvement following solution heat treatment. The fact that plastic deformation (*q*) was about 0.31 in the solution-treated condition means that the alloy reached 31% of its maximum quality index value (*Q*). The importance of *q* is that it shows how much a sample is away from its maximum possible ductility *q* = 1 and indicates that it would be possible to control the microstructure, for example by reducing the SDAS, or the porosity, or intermetallic level to enhance the alloy ductility and hence, the quality index, *Q*. When the ductility increases sharply from the as-cast to the solution heat treated condition, such changes can be related to the spheroidization of silicon particles and to the uniformity of the microstructure in the solution heat-treated condition, as shown in **Figure 6(a)**.

#### **Figure 6.**

Q*-charts following: (a) SHT, (b) aging at 155°C, (c) aging at 190°C, and (d) aging at 350°C. Legends in (b) apply for other charts. The curved lines indicate the passage from 2 to 10 to 100 h.*

*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

From the data presented in **Figure 6(b)**–**(d)**, it is evident that the change in crystallographic structure of Al2Cu phase from G-P zones (155°C) to a metastable phase (190°C) to a stable phase (350°C) is the main parameter controlling the alloy performance quality. As can be seen, at each aging temperature, all points fall within a narrow circle due the progress in the formation of the precipitated phase. The broken lines in these figures show the change in the *Q*-level as a function of aging temperature. The width of the circle deceased from 175 MPa (155°C) to 75 MPa (190°C) to 25 MPa (350°C), representing the hardening and softening behavior of the alloy as a function of the aging temperature and time [26]. Using aging times of 2 and 100 h as reference points, the *Q* , UTS and %El values are presented in **Table 5**. As can be seen, the *Q* values after 2 h are moreor-less same over such a large range of aging temperatures, due to the variation in both UTS and %El. However, aging for 100 h revealed highest value at 190°C compared to 155°C (under aging) and 350°C (over aging). The *Q* values for test bars aged at 350°C for 100 h is the same due to the balance between UTS and %El.

## **3.2 Stage II-alloy B**

The heat treatment procedures followed for the alloy B are listed in **Table 6**. The same treatments were applied for both 25°C and 250°C tensile testing.

**Figure 7** [27] shows the DSC heating curves of the alloys in the as-cast and SHT conditions, where three explicit peaks could be detected and coded 1, 2, 3. Considering the main parameter is the precipitation of Al2Cu phase particles, thus the height of peak number 1 following SHT compared to that in the as-cast condition plays a crucial role in controlling the alloy performance after aging. In addition, it is an indication of the effectiveness of the SHT process in dissolving the initial Al2Cu phase. In **Figure 7**, peak # 1 after solutionizing is more


#### **Table 5.**

Q*, UTS and %El values for alloy A after 2 and 100 h aging at different temperatures.*


#### **Table 6.**

*Heat treatment procedures and parameters applied to alloys investigated in stage II.*

#### **Figure 7.** *Portion of the DSC heating curves of as-cast and as-quenched alloy B samples [27].*

or less negligible due to dissolution of most of the Al2Cu phase, as shown in **Figure 4(b)**.

The principle phases seen in alloy B are demonstrated in the optical as well as backscattered (BSE) images displayed in **Figure 8(a)** and **(b)** [27], respectively. **Figure 8(a)** exhibits α-Al dendrites separated by eutectic silicon colonies. The phases observed in **Figure 8(b)** were identified using EDS analysis and reference to the results of Hernandez-Sandoval [28] and Garza-Elizondo [29]. Selective EDS spectra identifying these phases are displayed in **Figure 8(c)** through **Figure 8(d)**. The existence of Al2Cu phase in the block-like form may be attributed to the presence of Sr. in the alloy which leads to segregation of copper to localized areas [30]. The platelets of the Fe-rich β-Al5FeSi phase are easily recognized in the BSE image, surrounded by the blocky Al2Cu particles. The Mg-rich *Q*-phase (Al5Cu2Mg8Si6) is found growing out of the Al2Cu phase as seen in the BSE image. The absence of coarse Al3Zr precipitates [31] may be related to superheating that led to considerable dissolution of the Al3Zr phase from the master alloy during the melting process. As a result, the coarse Zr-containing phases are rarely detected since Al3Zr particles act as nucleation spots for these coarse phases. According to Garza-Elizondo [29], coarse Zr-rich particles may be nucleated on the undissolved Al3Zr particles provided by the master alloy, i.e., Al-15 wt.%Zr. In the present study, superheating the melt to 800°C would significantly reduce the numbers of Al3Zr particles in the matrix. The predicted fine zirconium trialuminide (Al3Zr) dispersoids that may be present on a nanoscale would require a high magnification BSE image to be detected.

**Figure 9(a)** shows a bright-field (BF) TEM image obtained in a T6-treated sample of alloy B with the electron beam parallel to the [001] zone axis. This figure shows a high density of uniformly distributed needle-like precipitates which are oriented along <110> family of directions and aligned along the [100] planes. The length of these precipitates ranges from 50 to 150 nm, close to the reported size of θ′-Al2Cu plates (50–100 nm long) [32, 33]. **Figure 9(b)** displays the associated selected area electron diffraction (SAED) pattern obtained from **Figure 9(a)**. The observable discrete diffraction maxima for the precipitates in SAED pattern indicate the presence of θ′-Al2Cu, where the streaks most probably result from the presence of fine S′-Al2CuMg particles. Computer simulation studies [34–37] on the S′-phase reflections show that they are hidden within the streaks of θ′.

*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

#### **Figure 8.**

*(a) Optical micrograph at 200× magnification, and (b) backscattered electron image of alloy B (354 + 0.3 wt.%Zr), obtained at a low cooling rate of 0.35°C/s, showing the different phases present in the alloy; (c–g) EDS spectra corresponding to Al2Cu, (Al,Si)3(Ti,Zr),* Q*-Al5Mg8Cu2Si6, (Al,Si)3Zr, and β-Al5FeSi phases observed in (b) [27].*

In the present work, **Figure 10(a)**, the addition of ~0.3 wt.%Zr to the 354-type Al-Si-Cu-Mg cast alloy in the as-cast condition improves the ambient-temperature (25°C) strength values of the Zr-free 354 alloy (alloy A), by ~26 MPa (UTS) and 40 MPa (YS), respectively. Following SHT, the UTS and ductility values remained almost constant at ~300 MPa and ~ 6.3%, respectively, while the yield strength increased by ~33 MPa compared to alloy A. It is believed that the improved strength

#### **Figure 9.**

*(a) Bright-field TEM image of alloy B in T6-treated condition, and (b) the selected area electron diffraction (SAED) pattern.*

#### **Figure 10.**

*(a) Ambient, and (b, c) high temperature tensile properties of alloy B.*

values of alloy B emphasizes the role of Zr addition in enhancing the ambient-temperature tensile properties through the formation of fine secondary strengthening precipitates (Al3Zr) as reported by many authors [14, 38–40]. The fact that UTS and YS in the T5 and T6 conditions are very close may be attributed to the strengthening effect of the fine dispersoids, which precipitate during the artificial aging stage of the T5, and T6 treatments as reported in **Figure 9**.

Tensile testing at 250°C, endured a significant softening due to the possible coarsening of the strengthening precipitates (Al2Cu) that existed during tensile testing at room temperature (**Figure 11**). In addition, the T5 heat treatment did

## *Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

not improve the elevated-temperature strength values of the as-cast alloys but reduced the alloy ductility by ~50%. However, application of the T6 heat treatment noticeably enhanced the strength values of the as-cast condition from about 175 to 225 MPa. Another parameter to consider is the effect of thermal stability. In the present work, some tensile samples were stabilized at 250°C (following the T5 and T6 aging treatment) for a lengthy period of time, i.e., 100 and 200 h. As can be seen from **Figure 10(c)**, the stabilized T5-treated alloy B samples exhibit better strength values (UTS and YS) than those obtained in the stabilized T6-treated condition. However, the ductility values obtained after stabilization of the T5-treated samples are dramatically lower in comparison. **Figure 11(a)** shows Al2Cu particle size and distribution in a T6 sample stabilized at 250°C for 200 h, whereas **Figure 11(b)** is the corresponding EDS spectrum.

A detailed investigation of the fracture surfaces of tensile bars of alloy B were examined in the T6-treated condition, before and after stabilization for 200 h at 250°C. The T6-temper treatment was selected due to its wide use in the automotive industry. The BSE image shown in **Figure 12(a)** [41] shows the fracture surface of the tensile-tested alloy in the T6-treated condition. The fracture surface

#### **Figure 11.**

*(a) Backscattered electron images showing the size and distribution of precipitates in the T6-treated B alloy after stabilization at 250°C for 200 h; (b) EDS spectrum corresponding to the rod-like particles in (a).*

#### **Figure 12.**

*Fracture surface of T6-treated alloy B: (a) BSE image showing a uniform dimple structure and cracked particles (arrowed), (b) EDS spectrum corresponding to the point of interest in (a), and (c) high magnification BSE image showing a cracked Al-Si-Ti-Zr particle (arrowed) [41].*

*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

#### **Figure 13.**

*Fracture surface of alloy B: (a, b) BSE images of T6-treated alloy after stabilization at 250°C for 200 h showing a coarse dimpled structure, coarsened precipitates and Alx(Zr,Ti)Si particles involved in the crack initiation process, and (c) corresponding EDS spectrum of the phase of interest as shown in (b) [41].*

has a dimpled-structure throughout, which indicates the ductile nature of the fracture mode. In addition, the BSE image exhibits the precipitation of Alx(Zr,Ti) Si compound, in the form of star-like shape, as confirmed by the associated EDS spectrum in **Figure 12(b)**. Also, cracks can be spotted in various particles of this compound, as indicated by the arrows. The higher magnification BSE image shown in **Figure 11(c)** reveals a cracked Alx(Zr,Ti)Si phase particle.

**Figure 13(a)** [41] shows the fracture surface of the T6-treated B alloy tested at 250°C after stabilization for 200 h at the testing temperature. The dimple structure is coarser compared to that before stabilization at 250°C. This observation would explain the improved ductility of the alloy due to the softening behavior associated with the prolonged elevated-temperature exposure at 250°C. Coarsened precipitates appear in the interiors of the dimples, as indicated by the oval contours in **Figure 13(a)**. The BSE image and the EDS spectrum shown in **Figure 13(b)** and **(c)**, respectively, confirm the presence of Alx(Zr,Ti)Si phase particles.

## **4. Conclusions**

Based on an analysis of the results presented in this article, the following conclusions may be made:


*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*


The Alx(Zr,Ti)1-xSi complex compound is observed with star-like and blocky morphologies, with cracks appearing in various particles of this compound. By increasing the stabilization time up to 200 h, coarser and deeper dimples are formed, highlighting the improved ductility of the alloy due to the softening behavior associated with the prolonged exposure at 250°C.

## **Author details**

Jacobo Hernandez-Sandoval1,2, Mohamed H. Abdelaziz1,3, Agnes M. Samuel1 , Herbert W. Doty4 and Fawzy H. Samuel1 \*

1 Département des Sciences appliquées, Université du Québec à Chicoutimi, Canada

2 Facultad de Ingeniería Mecánica y Eléctrica, Universidad Autónoma de Nuevo Leon, Mexico

3 Département PEC, Université Française d'Égypte, Le Caire, Egypt

4 General Motors Global Technology Center, Warren, MI, USA

\*Address all correspondence to: fhsamuel@uqac.ca

© 2020 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/ by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

## **References**

[1] Ibrahim MF, Samuel AM, Doty HW, Samuel FH. Effect of aging conditions on precipitation hardening in Al-Si-Mg and Al-Si-Cu-Mg alloys. International Journal of Metalcasting. 2017;**11**(2):274-286

[2] Hatch JE. Aluminum: Properties and Physical Metallurgy. USA: ASM International; 1984

[3] Samuel FH, Samuel A. Effect of heat treatment on the microstructure, tensile properties, and fracture behavior of permanent mold Al-10 Wt Pct Si-0.6 Wt Pct Mg/sic/10 P composite castings. Metallurgical and Materials Transactions A. 1994;**25**(10):2247-2263

[4] Ibrahim MF, Abdelaziz MH, Doty HW, Valtierra S, Samuel FH. Effect of microalloying elements on the heat treatment response and tensile properties of Al-Si-Mg alloys. In: Ares AE, editor. Solidification. Rijeka: IntechOpen; 2018. Ch. 01

[5] Hernandez-Sandoval J, Garza-Elizondo GH, Samuel AM, Valtierra S, Samuel FH. The ambient and high temperature deformation behavior of Al-Si-Cu-Mg alloy with minor Ti, Zr, Ni additions. Materials & Design. 2014;**58**:89-101

[6] Meyers C, Hinton K, Chou J-S. Towards the optimization of heat-treatment in aluminium alloys. Materials Science Forum. 1992;**72**:102-104

[7] Mohamed AMA, Samuel FH. A review on the heat treatment of Al-Si-Cu/Mg casting alloys. In: Czerwinsk F, editor. Heat Treatment; Conventional and Novel Applications. Rijeka: IntechOpen; 2012. Chapter 4

[8] Cayron C, Buffat P. Transmission electron microscopy study of the Β′ phase (Al–Mg–Si alloys) and Qc phase (Al–Cu–Mg–Si alloys): Ordering mechanism and crystallographic structure. Acta Materialia. 2000;**48**(10):2639-2653

[9] Matsuda K, Teguri D, Sato T, Ikeno S. EFTEM observation of Q'phase in Al-Mg-Si-Cu alloy. Materials Science Forum. 2002;**396**(4):947-951

[10] Hwang J, Banerjee R, Doty H, Kaufman M. The effect of Mg on the structure and properties of type 319 aluminum casting alloys. Acta Materialia. 2009;**57**(4):1308-1317

[11] Wang G, Sun Q, Feng L, Hui L, Jing C. Influence of Cu content on ageing behavior of AlSiMgCu cast alloys. Materials & Design. 2007;**28**(3):1001-1005

[12] Abdelaziz MH. Microstructural and mechanical characterization of transition elements-containing Al-Si-Cu-Mg alloys for elevated-temperature applications [PhD thesis]. Canada: UQAC; 2018

[13] Srinivasan D, Chattopadhyay K. Metastable phase evolution and hardness of nanocrystalline Al– Si–Zr alloys. Materials Science and Engineering A. 2001;**304**:534-539

[14] Knipling KE. Development of a nanoscale precipitation-strengthened creep-resistant aluminum alloy containing trialuminide precipitates [PhD thesis]. USA: Northwestern University; 2006

[15] Ragab K. The use of fluidized sand bed as an innovative technique for heat treating aluminum based castings [PhD thesis]. Canada: UQAC; 2012

[16] Mohamed A, Samuel A, Samuel F, Doty H. Influence of additives on the microstructure and tensile properties of near-eutectic Al–10.8% Si

*Effect of Zr Addition and Aging Treatment on the Tensile Properties of Al-Si-Cu-Mg Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.92814*

cast alloy. Materials & Design. 2009;**30**(10):3943-3957

[17] Drouzy M, Jacob S, Richard M. Interpretation of tensile results by means of quality index and probable yield strength-application to Al-Si7 Mg foundry alloys-France. International Journal of Cast Metals. 1980;**5**(2):43-50

[18] Jacob S. Quality index in prediction of properties of aluminum castings—A review. In: Transactions of the American Foundry Society and the One Hundred Fourth Annual Castings Congress. Vol. 208. Pittsburgh, Pennsylvania; 2000. pp. 811-818

[19] Cáceres C. A rationale for the quality index of Al-Si-Mg casting alloys. International Journal of Cast Metals Research. 2000;**12**(6):385-391

[20] Abdelaziz MH, Doty HW, Valtierra S, Samuel FH. Static versus dynamic thermal exposure of transition elements-containing Al-Si-Cu-Mg cast alloy. Materials Science and Engineering A. 2019;**739**:499-512

[21] Tahiri H, Mohamed SS, Doty HW, Valtierra S, Samuel FH. Effect of Sr–grain refining–Si interactions on the microstructural characteristics of Al–Si hypoeutectic alloys. International Journal of Metalcasting. 2018;**12**(2):343-361

[22] Zhu PY, Liu QY. Kinetics of granulation of discontinuous phase in eutectic structures. Materials Science and Technology. 1986;**2**:500-507

[23] Reif W, Dutkiewiewicz J, Ciach R. Effect of precipitates in Al-Si-Cu-Mg alloys. Materials Science and Engineering A. 1997;**234**:165-168

[24] Ammar HR, Samuel AM, Samuel FH. Porosity and the fatigue behavior of hvpoeutectic and hypereutectic aluminum-silicon casting alloys. International Journal of Fatigue. 2008;**30**(6):1024-1035

[25] Câceres CH, Taylor JA. Enhanced ductility in Al-Si-Cu-Mg casting alloys with high Si content. In: Tiryakiouglu M, Crepeau P, editors. Shape Casting: The John Campbell Symposium. California: TMS; 2005. pp. 245-254

[26] Toschi S. Optimization of A354 Al-Si-Cu-Mg alloy heat treatment: Effect on microstructure, hardness, and tensile properties of peak aged and overaged alloy. Metals. 2018;**8**:961. DOI: 10.3390/met8110961

[27] Abdelaziz MH, Elgallad EM, Doty HW, Valtierra S, Samuel FH. Melting and solidification characteristics of Zr-, Ni-, and Mn-containing 354-type Al-Si-Cu-Mg cast alloys. Philosophical Magazine. 2019;**99**(13):1633-1655

[28] Hernandez Sandoval J. Improving the Performance of 354 Type Alloy [PhD thesis]. Canada: Université du Québec à Chicoutimi; 2010

[29] Garza-Elizondo GH. Effect of Ni, Mn, Zr and Sc Additions on the Performance of Al-Si-Cu-Mg Alloys [PhD thesis]. Canada: UQAC; 2016

[30] Ibrahim MF, Samuel E, Samuel AM, Al-Ahmari AMA, Samuel FH. Metallurgical parameters controlling the microstructure and hardness of Al-Si-Cu-Mg base alloys. Materials & Design. 2011;**32**(4):2130-2142

[31] Priya P, Matthew JM, Krane JM, Johnson DR. Precipitation of Al3Zr dispersoids during homogenization of Al-Zn-Cu-Mg-Zr alloys. In: Light Metals. 2016. pp. 213-218

[32] Tavitas-Medrano F, Gruzleski J, Samuel F, Valtierra S, Doty H. Effect of Mg and Sr-modification on the mechanical properties of 319-type aluminum cast alloys subjected to artificial aging. Materials Science and Engineering A. 2008;**480**(1):356-364

[33] Reif W, Dutkiewicz J, Ciach R, Yu S, Krol J. Effect of ageing on the evolution of precipitates in AlSiCuMg alloys. Materials Science and Engineering A. 1997;**234**:165-168

[34] Purtee M. Aging effects of an aluminum-based 319 alloy. Foundry Management & Technology. 1998;**126**(10):42-44

[35] Takeda M, Ohkubo F, Shirai T, Fukui K. Precipitation behaviour of Ai-Mg-Si ternary alloys. In: Materials Science Forum, Vol. 217-222. Stafa-Zurich, Switzerland: Trans Tech Publications, Ltd.; 1996. pp. 815-820

[36] Elwazri A, Varano R, Siciliano F, Bai D, Yue S. Characterisation of precipitation of niobium carbide using carbon extraction replicas and thin foils by FESEM. Materials Science and Technology. 2006;**22**(5):537-541

[37] Kang H, Kida M, Miyahara H, Ogi K. Age-hardening behavior of Al-Si-Cu base cast alloys. Journal of Japan Foundry Engineering Society. 1997;**69**(10):828-834

[38] Shaha SK. Development and characterization of cast modified Al-Si-Cu-Mg alloys for heat resistant power train applications [PhD thesis]. Canada: Ryerson University; 2015

[39] Knipling KE, Dunand DC, Seidman DN. Nucleation and precipitation strengthening in dilute Al-Ti and Al-Zr alloys. Metallurgical and Materials Transactions A. 2007;**38**(10):2552-2563

[40] Knipling KE, Dunand DC, Seidman DN. Precipitation evolution in Al–Zr and Al–Zr–Ti alloys during isothermal aging at 375-425 C. Acta Materialia. 2008;**56**(1):114-127

[41] Abdelaziz MH, Elgallad EM, Samuel AM, Doty HW, Samuel FH. High-temperature tensile fractography of Zr-, Ni-, and Mn-containing Al-Si-Cu-Mg cast alloys. Advances in Materials Science and Engineering. 2020

## **Chapter 4**

## The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry

*Rafał Hubicki and Maria Richert*

## **Abstract**

This chapter describes and analyzes the 6xxx aluminum alloys used in the shape extrusion sector dedicated to automotive and construction industry. The division and application of 6xxx aluminum alloys are performed. The precipitation hardening of **6xxx** (Al-Mg-Si) alloys is presented as these alloys easily undergo deformation and present the potential for new kinds of alloys for high-speed extrusion. The mechanisms of strengthening are shown with the evolution of precipitation sequences. Also some examples of industry applications of 6xxx aluminum alloys are presented.

**Keywords:** aluminum alloys, extrusion, aging, microstructure

## **1. Introduction**

### **1.1 Characteristics of aluminum alloys**

Historically, aluminum was first produced in 1825 by reducing aluminum chloride with potassium amalgam. In 1886, Héroult and Hall discovered the possibility of producing aluminum by electrolysis. In 1895, aluminum was first used as a material for the church roof. With the increase in the production of aluminum, which occurred especially after the Second World War, the scope of its use expanded. Today, aluminum alloys are widely used in transport, mechanical engineering, electrical and energy industries, food industry, chemical industry, sports, aviation, transport, yacht and shipbuilding, and many other fields. Below, in **Table 1**, the grades of aluminum alloys used in industry are presented.


**Table 1.** *Industrial aluminum alloy groups.*

Aluminum alloys are highly ductile, so it is easy to make the desired structural elements, machine parts, and others from them.

By changing the content of alloying elements in aluminum alloys, the strength properties can be adjusted. A very effective factor influencing the strengthening of aluminum alloys is heat treatment—supersaturation and aging. This treatment is possible for Al-Mg-Si 6xxx series alloys, which show a variable solubility in the solid state. Similar possibilities are also available for the 2xxx, Al-Cu-Mg, and 7xxx, Al-Zn-Mg-Cu, alloys (**Figures 1** and **2**). However, the most popular alloy for profile extrusion is 6xxx series.

During the extrusion process, aluminum alloys are placed in the extrusion press container (cylinder) and pressed by a pressing ram (or stem—via a dummy block or pressure plate). The metal flows out through the hole in the die, which gives the shape to the extruded profile (**Figure 3**).

The state of stress in most of the plasticized area is a three-axial nonuniform compression. It is therefore possible to make large plastic deformations without affecting the consistency of the material (maximum elongation coefficients are about 300, average—about 50). This is the main advantage of extrusion processes. Large deformations require high forces. The main limitation to the scale of deformation that can be obtained in one extrusion operation is not the phenomenon of material decohesion (as in many other processes) but the strength of the tools.

When the extruded section leaves the tool, it is cooled with water or air and then drawn, still in the malleable state. This removes the stresses accumulated in the aluminum alloy and at the same time allows to achieve the expected and correct profile dimensions. The profiles are then cut and obtain ultimate strength by hot or cold hardening.

#### **Figure 1.**

*Isothermal section of ternary Al-Cu-Mg phase diagram at 400°C, h ¼ CuAl 2, S¼ MgCuAl 2, and T ¼ (Cu 1-x Al x) 49 Mg 32. Along the dashed red line, the atomic fraction is constant, 94Al-6 Mg. And at the red star, the composition point is 90Al-4Cu-6 Mg, namely, the average composition of the eutectic region, shown with the dashed white box.*

*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

#### **Figure 2.**

*Phase diagram Al-Cu-Mg-Zn: (a) polythermal diagram, (b) distribution of phase fields in the solid state in the aluminum corner, and (c) single-phase domains [1].*

**Figure 3.** *Extrusion diagram and extrusion press for aluminum alloys.*

In the case of large strains, hot extrusion is used, because during cold extrusion the forces are so high that the tools do not withstand the loads. Cold extrusion for large deformations can only be performed for soft materials (e.g., pure aluminum).

## **2. 6xxx series aluminum alloys for extrusion of shapes**

The 6xxx series alloys are the most commonly used and their global consumption is the largest. The worldwide demand for aluminum is around 29 million tons per year. About 22 million tons is new aluminum and 7 million tons is recycled aluminum scrap. The use of recycled aluminum is economically and environmentally compelling. It takes 14,000 kWh to produce 1 ton of new aluminum. Conversely, it takes only 5% of this energy to remelt and recycle 1 ton of aluminum. There is no difference in the quality between virgin and recycled aluminum alloys [2].

Alloys of the 6xxx series are heat-treated as they show a variable solubility in the solid state. **Figure 4** shows a pseudo-binary system Al-Mg2Si. The variable solid solubility curve allows for heat treatment of 6xxx series alloys. After plastic working (e.g., extrusion), the supersaturation and artificial aging of these alloys are applied. After supersaturation, which consists in rapid cooling of the alloy, it is then heated to an appropriate temperature in order to precipitate hardening phases.

The phase sequence in this alloy is as follows:

$$\begin{array}{l}\text{super-saturated solid solution (SSSS)}\\-\text{atomic clusters} - \text{GP zones } - \text{ $\boldsymbol{\beta}$ '} - \text{ $\boldsymbol{\beta}$ '} / \text{ $\mathbf{\dot{O}}$ '} - \text{ $\boldsymbol{\beta}$ '} \,\text{S} \end{array} \tag{1}$$

**Figure 4.** *Calculated equilibrium phase diagram of Al-Mg2Si pseudo-binary alloys with excess Mg.*


*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

### **Table 2.**

*Precipitations in 6xxx series alloys.*

**Table 2** describes the characteristics of the precipitating phases [12].

Si plates appear in silicon-rich alloys. All phases that show the shape of needles are consistent with the direction <100> of the aluminum matrix.

The extrusion temperature of Al-Mg-Si alloys ranges between 400°C and 500°C, which causes that all phases are dissolved in the alloy to form a solid solution. Therefore, in order to prevent the precipitation of an incoherent phase of Mg2Si, a rapid quenching of the alloy is required after extrusion.

Rapid quenching maintains the supersaturation of the alloy's solid solution caused by Mg and Si and a high concentration of vacancies, also resulting from rapid cooling. The supersaturation is followed by the aging process (**Figure 5**). It is carried out at temperatures from 165°C to 185°C until maximum hardness is reached—state T6.

Hardening achieved as a result of aging depends on the size and density of the precipitations and the volume of metastable phases obtained.

Isothermal heating at 175°C induces the precipitation of β-phase hardening the alloy. The β″-phase is the most effective hardening phase in Al-Mg-Si alloys and is

**Figure 5.** *Sketch of precipitation hardening phases.*

formed at temperatures between 125°C and 200°C. The β″-phase composition is Mg5Si6; therefore, an appropriate Mg/Si ratio of 5/6 should be maintained during the casting process. The β-phase (Mg2Si) is not formed below 200°C, so it cannot be present in alloys aged below this temperature.

Marioara et al. [12] studied phase β″ precipitation at different Mg/Si ratios. The research showed that a higher Si content in the alloy promotes the formation of a large number of fine GP zones in a shorter time than alloys with a lower Si content. After annealing for 3 h at 175° C, there was a sharp peak of hardness associated with the appearance of GP zones in these alloys, and after 17 h, a wider peak of hardness associated with the occurrence of phase β″ precipitation appeared. Alloys richer in Mg had thicker elements of microstructure than alloys with an increased Si content. They also contained less U2 precipitations but were richer in β′ precipitations. The studies carried out revealed a strong influence of Si, the content of which controls the process of phase precipitation through the formation of precipitation clusters in the initial stages of annealing.

The great interest in Al-Mg-Si alloys is related to their application in the automotive industry. Chakrabarti et al. [13] conducted a research proving that the strengthening phase involved in the Al-Mg-Si ternary alloys is the metastable β″ phase [14, 15].

Kuroda et al. studied the effect of small amounts of Ni, Co, and V on hardness after age hardening of the AlMgSi base alloy. It was found that the additives increase the hardness, with the highest value of hardness occurring after the introduction of Ni and two-stage aging [16].

The influence of Cu addition on the precipitation process in AlMgSi alloy was studied by Zandbergen et al. [17]. The occurrence of Cu in all investigated precipitations of the hardening phases of the alloy was determined.

On the other hand, the influence of Zn and Ag additives on the precipitation process in AlMgSi alloys was analyzed in the work of Saito et al. [18]. The researchers found a weak influence of Zn on the precipitation sequence of hardening phases. They found that Zn is built into the structure of precipitations replacing some elements in the network of precipitated phases.

A new alloy based on AlMgSi, with the addition of Mn that did not require aging after extrusion, was presented by Lee et al. [19]. Mn creates precipitations of 0.05–0.5 μm. The alloy containing 1 wt.% Mn had high strength properties and good ductility in comparison with the commercial 6N01 alloy. Moreover, it showed higher fatigue resistance.

Chen et al. studied the process of phase precipitation in Al-Mg-Si alloys [20]. They showed that the formation of Mg5Si6, hardening AlMgSi alloy, is preceded by the precipitation of Mg2Si2,6Al6,4 and more precisely Mg2Si2Al7. The precipitation of this phase is dependent on the temperature of the aging of the alloy.

Mechanisms of hardening with β″ particles, characterized by coherence with aluminum network, were studied by Ringdalen et al. [21]. The study concerned the alloy 6060 and included the interaction of the dislocation lattice with the precipitations.

The AlMgSi alloy, due to the possibility of heat treatment, has a high potential for modeling its structures and properties. Analyzing the process of supersaturation of this alloy and subsequent aging, it is possible to obtain the assumed properties. Hardening phases in the AlMgSi alloy nucleate heterogeneously at the grain boundaries and iron precipitations, which are always present in AlMgSi alloys

**Figure 6.** *Hardening phases in AlMgSi alloy.*

(**Figure 6**). Precipitations of β″ , the hardening phase of AlMgSi alloy, can be found at **Figure 7**. Their growth to larger sizes is associated with the formation of Mg2Si phase. Large Mg2Si particles are not an obstacle to displaced dislocations and do not increase the strength properties of the alloy.

An example of research on the heat treatment of AlMgSi alloys is the article by Richert et al. [22] concerning the new, multistage aging of the 6xxx series alloys. The studies of Ryen et al. [23] showed the presence in alloys AA6063 and AA6068, after aging precipitations as B′, β″, U2, and U1, β′ and some unidentifiable crystal structures. Detailed studies of the precipitations in aluminum alloys are also presented in the work by Andersen et al. [24]. From this work, data were taken from the composition of the precipitations formed during aging in AlMgSi alloys:

β-Mg2Si, U1-MgAl2Si2 'A' β″-Mg4(AlxMg1-X)Si4 U2-Mg4Al4Si4 β′-Mg9Si5 B′ ~ Mg48Al16Si36 GP-Mg4(AlxMg1-x)Si4

The author found that the early stages in Al-Mg-Si and Al-Mg-Cu systems have isostructural GP zones being 1D strings along <100>Al, identical to the eye-like units of the β″-phase, where Cu can replace Si. Calculations show GP zones can take different compositions, the most stable being Mg4Si4Mg and Mg4Cu4Al, with Mg and Al as central interstitial columns. Solute clusters for the GP zones are likely short defect-free needles using a vacancy to produce a central interstitial column. The needle-shaped β″-phase is the most important hardening precipitate in the AA6xxx system.

**Figure 7.** *Nanometric precipitations in AlMgSi alloy [25].*

## **3. High-speed aluminum alloy**

## **3.1 Studies on the production of a high-speed extrusion alloy**

## *3.1.1 Research methodology*

In this work, tests were carried out on 6060 alloy with three different chemical compositions, falling within the definition of the standard EN-AW 755-2 (**Table 3**).

Alloy 1 with the lowest Si and Mg contents, Alloy 2 with medium Si and Mg contents, and Alloy 3 with the highest Si and Mg contents were prepared as follows (**Figure 8**):

	- Alloy 1 (low alloying content)
	- Alloy 2 (with medium content of alloying elements—standard composition)
	- Alloy 3 (high alloying content)
	- From each alloy, three billets were extruded with the standard speed of the extrusion ram for this shape, and three billets with the extrusion ram speed increased by 20%.
	- The temperatures of the extrusion billets and the press container were the same for all billets.
	- All extruded sections were saturated on the press runout table and then artificially aged in the furnace to T66.
	- Three samples of each extruded section were taken for each hole to check the strength properties and material hardness of the sections.


*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

**Table 3.**

*Chemical composition of billets tested.*

#### **Figure 8.**

*Aluminum alloy billets before loading onto the loading ramp of the press [26].*

The tests of mechanical properties of extruded sections were carried out on the Zwick Z300 fatigue tester.

1.In the second stage, only Alloy 1 was tested as promising the highest ductility and thus the possibility to use the highest extrusion rate for the manufactured sections.

The following tests were carried out:


#### **Table 4.**

*Aging curves of Al-Mg-Si alloy with the lowest composition of alloying elements, second stage of research.*


## *3.1.2 Comparative results of tested alloys*

The research revealed that after casting and homogenization, Alloy 1 has a correct structure with equiaxed grains and broken lattices of iron precipitations (**Figure 9**).

**Figure 10** shows the yield point (R0.2) and tensile strength (Rm) of extruded profiles made of Alloy 1 with the lowest alloying content and other alloy compositions with higher content of alloying elements, i.e., Alloy 2 and Alloy 3. According to the requirements of the standard, the extruded section should have minimum R0.2 = 160 MPa and Rm = 215 MPa. The shape made of Alloy 2 with the average content of alloying elements and Alloy 3 with the highest content of alloying elements obtained the property level required by the standard EN755-2 (for products made of alloy 6060, in the T66 state). The properties of Alloy 3 are too high. And the shape with the lowest level of alloying elements failed to obtain such a level.

#### **Figure 9.**

*The structure of Alloy 1 billet with the lowest content of Mg and Si alloying elements.*

#### **Figure 10.**

*Yield point and tensile strength of extruded sections from billets with three different contents of the main alloying elements.*

**Figure 11** illustrates the plasticity of the alloys tested. The highest value of elongation was achieved by the section with the lowest content of alloying elements. This result indicates that the alloy potentially has a plasticity level allowing to use higher extrusion rates in relation to the other alloying components.

The delta (**Figure 11**), i.e., the difference between R0.2 and Rm, shown in the diagram (**Figure 11**) illustrates the "yield strength reserve" and confirms the favorable deformation conditions for the alloy with the lowest level of alloying elements with respect to the potential of utilizing its higher plasticity.

The yield point and tensile strength level of Alloy 1 must be achieved by means of a new, finely developed heat treatment.

The next figure (**Figure 12**) shows the elongation of the extruded sections and their hardness. The obtained result confirms that the alloy with the lowest content of alloying elements (Alloy 1) has the highest yield strength reserve.

## *3.1.3 Studies on the effects of the extrusion rate on the properties, structures, and phase compositions of the alloy with the lowest content of alloying elements*

The potential possibility of using the cheapest Alloy 1 with the lowest content of alloying components while optimizing the extrusion rate for the most efficient *The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

#### **Figure 11.**

*Elongation and yield strength reserve of extruded sections from billets with three different contents of the main alloying elements.*

**Figure 12.**

process created grounds for research into heat treatment optimization to achieve the properties required by the standard. Alloy 1 has been tested for its structures, properties, and phase compositions.

**Table 5** shows the macrostructure and microstructure of extruded Alloy 1 shapes with respect to the increasing extrusion rate. It is characterized by the equiaxed grains of medium size (medium chord parameter) in the range from 46 μm to 54 μm. Homogenization of the structure and grain size with a maximum grain size difference of 15% results from the hot working, activating the processes of structure recovery during extrusion.

Results of phase composition tests of extruded shapes are presented in **Figure 13**.

There is a final slight decrease in the content of the precipitated phases at the highest extrusion rate applied standard value +60%. At the standard rate +20%, the highest phase content of Mg2Si and Mg5Si6 was recorded. The obtained result sets the direction for actions aimed at obtaining sections with assumed properties at the highest efficiency of the extrusion process. An increase in the extrusion rate above 20%, without changing the subsequent heat treatment applied after extrusion to the


**Table 5.** *Structure of Alloy 1 sections depending on the extrusion rate.*

T6 state, is an optimal variant ensuring the highest efficiency of the process. On the other hand, if the applied heat treatment allowed for a significant increase in the hardening phase, then higher extrusion rates could be applied.

## *3.1.4 Effect of heat treatment on the properties of extruded sections made of Alloy 1*

The alloy meets the criteria of a high-speed extrusion alloy as the tests resulted in a significant increase in the extrusion rate of sections in relation to the rate currently used. However, due to the low content of Mg and Si alloying elements, the

*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

#### **Figure 13.**

*Phase changes in extruded sections.*


#### **Table 6.**

*Final results of strength properties after heat treatment for Alloy 1: stage 2.*

classical heat treatment did not allow to obtain the recommended level of properties according to the standard. Therefore, it was necessary to carry out a number of studies to develop a new heat treatment that would ensure that the yield point, tensile strength, and hardness value of the standard were achieved. **Table 6** contains data on the parameters of the applied heat treatment variants.

The analysis of the obtained data showed that in variant 19 of heat treatment, the yield point reached the level assumed in the standard. The obtained result is a great technological achievement and promises a better economic result.

The research has shown that it is possible to produce a cheaper alloy (low content of base alloying elements), which allows to use a 60% higher extrusion rate and, using an appropriate heat treatment, achieve a comparable level of properties to alloys containing higher levels of Mg and Si elements.

The higher ductility of the new high-speed extrusion alloy enables the production of a greater tonnage of products per unit of time than the currently used alloys, which translates into increased production efficiency.

## **4. Summary: 6xxx aluminum alloys for a high-speed extrusion**

At the beginning of the twenty-first century, commercial concerns involved in the production and sale of extrusion sections made of aluminum alloys started to show interest in increasing the productivity of the extrusion process through a radical change in the technology. The subgrades of alloy 6060 with a reduced content of alloying elements, called lean alloys or high-speed alloys, began to form. They were characterized by an increased plasticity in relation to classical alloy 6060, as a result of which the speed of the process increased by 20% in comparison with alloys with classical chemical composition. In order to obtain certain strength properties, it was necessary to use unconventional aging cycles of the material. Because the alloys required an increase in extrusion rate, lean alloys were used primarily for extrusion of sections with relatively uncomplicated cross sections. These types of sections were mainly used in the construction industry, so they were sometimes called building alloys. By 2007, thanks, among other things, to the exchange of workers between the companies, the production technology of building alloys became widely known, and the hydro concern started to produce billets from these alloys for sale to the general public, designing and producing new generations of alloys for fast extrusion at the same time.

Designing alloys for high-speed extrusion requires a deep knowledge of the hardening mechanisms and plastic deformation processes. During hot deformation, which takes place in extrusion, it is important to know the effect of temperature on the extrusion processes. The processes of structure recovery during hot deformation are dynamic recrystallization and dynamic recovery. In aluminum and its alloys, which exhibit high stacking-fault energy, dynamic recovery processes take place.

All effects in the extrusion process are temperature and speed controlled. It is shown that the mechanism of deformation is either dynamic recovery or dynamic recrystallization, according to whether the alloy is of high or low stacking-fault energy. This leads to extrudate structures which contain either subgrains, ideally with no static recrystallization occurring, or grains formed by successive dynamic, metadynamic, and static recrystallization. It is shown that either type of structure is related to, and can be controlled by, the prevailing Zener-Hollomon (Z-H) parameter *Z.* There is thus also a relationship between properties and *Z* [27].

The Zener-Hollomon parameter is given by equation:

$$Z = \dot{\mathbf{e}} \exp(\mathbf{Q}/\mathbf{RT})\tag{2}$$

where e is the strain rate, Q is the activation energy, R is the gas constant, and T is the temperature.

The Zener-Hollomon parameter decreases with a decreasing strain rate and an increasing temperature. In the actual industrial production, to avoid the occurrence of cracks, low Z conditions (lower strain rates and higher temperatures) are usually preferred; however, when the deformation rate is too low, the temperature drop is serious, which is harmful for deformation [28].

In case of high-speed extrusion alloys Z-H parameter should increase, the strain rate will increase and temperature should be lowered. The decrease of subgrain/ grain size should be expected.

Deformation temperature and strain rate are important factors controlling the hot deformation flow stress. The experimental results show that the dynamic softening is accelerated with the increase of deformation temperature and decrease of strain rate. Strain effects on flow stress and hence on extrusion pressure are predominant for hot extrusion (due to strain rate sensitivity). Therefore, it is rather difficult to predict the extrusion force in hot extrusion. We can estimate the strain rate at any location x in the billet from the geometrical considerations. Let a cylindrical billet have an initial radius of Ro and extruded radius of Rf. be semi-cone angle of the die.

We can formulate the strain rate at any location x from the entry of die as

*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

$$\dot{\varepsilon} = \frac{d\varepsilon}{dt} = -2 \frac{\mathbf{v}\_0 \mathbf{R}\_0^2}{\left(\mathbf{R}\_o - \mathbf{x} \tan \alpha\right)^3} \tan \alpha$$

The average strain rate undergone by a billet is given by

$$\csc \frac{\mathsf{G} \mathbf{v}\_0 \mathbf{D}\_0^\natural \mathsf{tan} \; \alpha}{\left(\mathbf{D}\_0^\natural \mathsf{-} \mathbf{D}\_\mathbf{f}^\natural\right)} - \ln \mathbf{R} \tag{4}$$

where v0 is the velocity of ram.

It is generally believed that the average grain size (*D*A) decreases with increasing strain rate and decreasing temperature and is independent of initial grain size and accumulated strain.

The effects of temperature and strain rate can be expressed in terms of the second-order function of the Zener-Hollomon parameter, *Z*, in an exponent-type function of temperature and strain rate, as follows [29]:

$$D\_{\Lambda} = B\_{\circ} + B\_{\circ} \ln Z + B\_{\circ} \left(\ln Z\right)^{2} \tag{5}$$

where *B*1 and *B*2 are polynomial coefficients. Generally, the grain size of asextruded material decreases with the increase of the Zener-Hollomon parameter. The fact that the Z-H parameter determines the microstructural evolution reveals that the foundation of an ultrafine-grained structure is formed dynamically and the process is closely related to the thermal activation process during and after the deformation.

In general, dynamic recrystallization does not seem to occur in aluminum alloys with Mg below 4%. The high recovery processes develop in such aluminum alloys—with subgrain lattice formation.

The size of stable subgrains that result from dynamic recovery in aluminum alloys depends upon hot working conditions. McQeen et al. [30] expressed this by equation:

$$\mathbf{d}\_s = \mathbf{a} + \mathbf{b}\log\mathbf{Z} \tag{6}$$

where ds is the subgrain size, Z is the Zener-Hollomon parameter, and a and b materials are constant.

The subgrain size usually reaches a limiting value of a few microns in aluminum alloys. With continued deformation, there is usually no further decrease in the subgrain size.

For hot extrusion, the extrusion pressure p is directly proportional to stain rate. As strain rate increases, the extrusion pressure also increases, almost linearly. As ram speed increases, the extrusion pressure also increases, due to increasing strain rate. However, the extrusion pressure is reduced with increased working temperature in hot extrusion (**Figure 14**). The extrusion speed has some limit depending on billet temperature. It means that the technological application of high-speed alloys requires some experimentally developed extrusion parameters.

**Figure 14.** *Temperature dependence to extrusion speed.*

In the analysis of hardening of polycrystalline materials, the component (σy), described by the model of the Hall-Petch (H-P) equation, acquires a particular importance as it allows to determine the value of lower yield point as a function of grain size:

$$
\mathbf{\dot{o}}\_{\mathbf{y}} = \mathbf{\dot{o}}\_{\mathbf{o}} + \mathbf{k}\_{\mathbf{y}} \mathbf{d}^{-\mathbf{a}/\mathbf{z}} \tag{7}
$$

where σy: normal stress corresponding to the yield point; σ0: stress of internal friction of the mobile dislocations; ky: slope factor, characterizing the resistance of grain boundaries to dislocation motion; d: diameter of grain or subgrain.

For this reason, the selection of conditions of extrusion rate and temperature determining the grain size is of great importance in the technology of extrusion of aluminum alloys.

At present, high-speed extrusion alloys are a desired alternative to conventional alloys due to the prospect of higher tonnage throughput per unit of time. The expected effect of this action is an increase in the company's profit. In the international literature, it is difficult to find data on the chemical composition of alloys for a high-speed extrusion because they are confidential information, which results from the well-understood interest of the company. There is also a noticeable lack of wider interest of scientists in this subject, who usually strive to disseminate the results, as it is contrary to companies' confidentiality requirements. Therefore, this type of research is usually conducted in a narrow circle of professionals closely cooperating with the industry. Among the literature items and reports that appear on the websites, there is information about progress in the production of new aluminum alloys for specific purposes, e.g., for the automotive, construction, high-speed rail, and other industries. The purpose of the alloys determines their production and properties.

The research conducted on fast extrusion alloys is a promising direction for the development of section production and provides prospects for improving extrusion technology. Great possibilities of increasing the productivity of the extrusion process connected with the implementation of these new quality aluminum alloys are also an element of the struggle for the market and increase in the productivity of companies.

*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

## **Author details**

Rafał Hubicki1 \* and Maria Richert<sup>2</sup>

1 Grupa Kęty S.A., Poland

2 AGH University of Science and Technology, Krakow, Poland

\*Address all correspondence to: h\_rafal@hotmail.com

© 2020 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/ by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

## **References**

[1] Butu M, Moldovan P, Marcu FD, Ungureanu I. Thermodynamics of in situ production of aluminium matrix composites comparative analysis. Materiale Plastice. 2016;**53**(3):428-433

[2] Available from: https://www.azom. com/article.aspx?ArticleID=2863

[3] Vissers R, Van Huis MA, Jansen J, Mariora CD, Andersen SJ. The crystal structure of the β′ phase in Al–Mg–Si alloys. Acta Materialia. 2007;**55**(11):3815-3823

[4] Marioara CD, Andersen SJ, Jansen J, Zandbergen HW. Atomic model for GP-zones in a 6082 Al–Mg–Si system. Acta Materialia. 2001;**49**:321-328

[5] Andersen SJ, Zandbergen HW, Jansen J, Traeholt C, Tundal U, Reiso O. The crystal structure of the β″ phase in Al–Mg–Si alloys. Acta Materialia. 1998;**46**(9):3283-3298

[6] Derlet PM, Andersen SJ, Marioara CD, Froseth A. A firstprinciples study of the β″ phase in Al-Mg-Si alloys. Journal of Physics: Condensed Matter. 2002;**14**:4011-4024

[7] Froseth A. PhD thesis. Trondheim: Norwegian University of Science and Technology; 2003. pp. 63-111

[8] Andersen SJ, Marioara CD, Froseth A, Vissers R, Zandbergen HW. Crystal structure of the orthorhombic U2-Al4Mg4Si4 precipitate in the Al– Mg–Si alloy system and its relation to the β′ and β″ phases. Materials Science Journal Engineering A. 2005;**390**(1-2):127-138

[9] Matsuda K, Sakaguchi Y, Miyata Y, Uetani Y, Sato T, Kamio A, et al. Precipitation sequence of various kinds of metastable phases in Al-1.0mass% Mg2Si-0.4mass% Si

alloy. Journal of Materials Science. 2000;**35**:179-189

[10] Madsuda K, Ikeno S, Sato T, Kamio A. A metastable phase having the orthorhombic crystal lattice in an Al-1.0mass% Mg2Si-0.4mass% Si alloy. Scripta Materilia. 1996;**34**:1797-1802

[11] Froseth A, Hoier R, Derlet PM, Andersen SJ, Marioara CD. Bonding in MgSi and Al-Mg-Si compounds relevant to Al-Mg-Si alloys. Physical Reviews B. 2003;**67**:224-106

[12] Marioara CD, Andersen SJ, Zandbergen HW, Holmestad R. The influence of alloy composition on precipitates of the Al-Mg-Si system. Metallurgical and Materials Transactions A. 2005;**36A**:691-702

[13] Chakrabarti DJ, Peng Y, Laughlin DE. Precipitation in Al-Mg-Si alloys with Cu additions and the role of Q' and related phases. In: Materials Science Forum. Vol. 396-402. USA: Alcoa Technical Center; 2002. pp. 857-862

[14] Edwards GA, Stiller K, Dunlop GL, Couper MJ. The Precipitation Sequence in Al-Mg-Si Alloys. Acta Materialia. 1994;**46**:3893-3904

[15] Chakrabarti DJ, Cheong BK, Laughlin DE. Automotive Alloys. USA: TMS; 1998. p. 27

[16] Kuroda Y, Yoshino D, Lee SW, Ikeno S, Matsuda K. Microstructure of small amount of TM added Al-Mg-Si alloys with two step ageing. In: Proceedings of 11th Polish-Japanese Joint Seminar on Micro and Nano Analysis, Gniew, September 11-14. 2016. DOI: 10.12693/APhysPolA.131.1373

[17] Zandbergen MW, Cerezo A, Smith GDW. Study of precipitation in Al-Mg-Si alloys by atom probe tomography. II. Influence of

*The High-Speed 6xxx Aluminum Alloys in Shape Extrusion Industry DOI: http://dx.doi.org/10.5772/intechopen.93239*

Cu additions. Acta Materialia. 2015;**101**:149-158

[18] Takeshi Saito S, Wenner A, Osmundsen CD, Marioara SJ, Andersen J, Royset W, et al. The effect of Zn on precipitation in Al-Mg-Si alloys. Philosophical Magazine. 2014;**94**:2410-2425. DOI: 10.1080/14786435.2014.913819

[19] Lee DH, Park JH, Nam SW. Enhancement of mechanical properties of Al-Mg-Si alloys by means of manganese dispersoids. Materials Science and Technology. 2013;**15**(4):450-455

[20] Chen JH, Constan E, van Huis MA, Xu Q, Zandbergen HW. Atomic pillarbased nanoprecipitates strengthen AlMgSi alloys. Science. 2006;**312**:416- 419. DOI: 10.1126/science.1124199

[21] Ringdalen I, Wenner S, Friis J, Marian J. Dislocation dynamics study of precipitate hardening in Al-Mg-Si alloys with input from experimental characterization. MRS Communication. 2017;**7**(3):626-633. DOI: 10.1557/ mrc.2017.78

[22] Richert J, Mroczkowski M, Gellner J. Nowy sposób starzenia stopów AlMgSi wyciskanych w stanie T5. Rudy i Metale Nieżelazne. 2006;**51**(12):747-753

[23] Ryen Ø, Holmedal B, Marthinsen K, Furu T. Precipitation, strength and work hardening of age hardened aluminium alloys. Materials Science and Engineering. 2015;**89**:012013. DOI: 10.1088/1757-899X/89/1/012013

[24] Andersen SJ, Marioara CD, Friis J, Wenner S, Holmestad R. Precipitates in aluminium alloys. Advances in Physics: X. 2018;**3**(1):1479984. DOI: 10.1080/23746149.2018.1479984

[25] Available from: https://www.bing. com/images/search?view=detailV2&cci d=V60%2B5rNV&id=019EC72102CAE6 7AB922781ADE5BCA40109097CB&thid =OIP.V60-5rNVOMKKmWU8zuw79gHa Ff&q=mg2si+precipitations+in+6060+a lloy&simid=608017416757841339&sele ctedindex=1&mode=overlay&first=1

[26] Grupa Kęty S.A. Photos by authors.

[27] Sheppard T. Temperature and speed effects in hot extrusion of aluminium alloys. Metals Technology. 1981;**8**(1):130-141. DOI: 10.1179/030716981803276009

[28] Li J, Liu B, Wang Y, Shan T, Liu Y, Lu X. A study on the Zener-Hollomon parameter and fracture toughness of an Nb-particles-toughened TiAl-Nb alloy. Metals. 2018;**8**(4):287. DOI: 10.3390/ met8040287

[29] Quan G-z, Wang Y, Liu Y-Y, Zhou J. Effect of temperatures and strain rates on the average size of grains refined by dynamic recrystallization for as-extruded 42CrMo steel. Materials Research. 2013;**16**(5):1092-1105. DOI: 10.1590/S1516-14392013005000091

[30] McQeen HJ, Celliers OC. Application of hot workability studies to extrusion processing: Part II. Microstructure development and extrusion of Al, Al-Mg and Al-Mg-Mn alloys. Canadian Metallurgical Quarterly. 1996;**35**(4):305-319

## **Chapter 5**
