The effect of spheroidisation heat treatment on the creep resistance of a cast AlSi12CuMgNi piston alloy

Materials Science & Engineering A 598 (2014) 147–153

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The effect of spheroidisation heat treatment on the creep resistance of a cast AlSi12CuMgNi piston alloy R. Fernández-Gutiérrez n, G.C. Requena Institute of Materials Science and Technology, Vienna University of Technology, Karlsplatz 13/308, A-1040 Vienna, Austria

art ic l e i nf o

a b s t r a c t

Article history: Received 14 August 2013 Received in revised form 5 December 2013 Accepted 28 December 2013 Available online 8 January 2014

The isothermal creep resistance of an overaged AlSi12CuMgNi alloy was studied at 300 1C as a function of solution heat treatment time at 480 1C. The microstructure of the alloy in as cast condition is formed by a three dimensional network of Si lamellae and intermetallic compounds embedded in an age-hardenable α-Al matrix. The solution heat treatment provokes the partial dissolution of Mg-rich intermetallic compounds and the spheroidisation of eutectic Si. These microstructural changes reduce the creep resistance of the alloy with increasing solution heat treatment time. The determination of the load transfer loss during solution heat treatment and an analogy with a modified shear-lag model for metal matrix composites indicate that the largest reduction in creep resistance occurs in the first 20 min–1 h of solution heat treatment. & 2014 Elsevier B.V. All rights reserved.

Keywords: Aluminium alloys Casting Mechanical characterisation Creep

1. Introduction Cast Al–Si alloys are widely used for engine blocks, pistons and cylinder heads due to their high fluidity and relatively high strength-to-weight ratio [1]. The creep resistance and high temperature strength of these alloys are determined by the strength of the microstructural components, their spatial distribution, volume fraction, morphology, size and interconnectivity. Particularly, rigid phases such as eutectic Si and intermetallic compounds can carry part of the load transferred from the matrix if they are present in sufficiently large volume fractions (e.g. [2–7]). The solidification of Al–Si alloys during industrial casting conditions results in the formation of highly interconnected three dimensional networks of eutectic Si [8]. Solution heat treatment (ST) of these alloys produces a gradual spheroidisation of the eutectic Si that provokes the disintegration of this network and the rounding of the eutectic Si. As a consequence, the load bearing capability of Si is reduced and the strength of the alloy decreases [2]. The addition of transition elements such us Cu, Mg or Ni to cast Al–Si alloys has a positive effect on the high temperature strength (e.g. [9–11]) through the formation of stable aluminides. The latter are highly contiguous with the eutectic Si [12] and partially hinder the spheroidisation of Si. This effect maintains the load bearing capability and the strength of the alloy [13].

The creep resistance of cast Al–Si alloys has received some attention in the past, usually as reference matrices of discontinuously reinforced metal matrix composites (MMC) [14–18]. However, no attention has been paid to the effects that solution heat treatment may have on the internal structure of the alloys and their influence on creep strength. The present work investigates these effects on a cast AlSi12CuMgNi alloy to help understand the process of load partition between the age hardenable α-Al matrix, eutectic Si and intermetallic compounds. It also sheds light on previous apparently controversial results [14–18]. 2. Experimental 2.1. Material An AlSi12CuMgNi alloy with the chemical composition shown in Table 1 was produced by squeeze casting at the Institute of Materials Science and Technology of the TU Clausthal [19]. The alloy was studied in as cast (AC) condition and after solution heat treatment at 480 1C over 20 min, 4 h and 64 h periods and followed by air cooling. Afterwards, artificial overageing was carried out for 2 h at 300 1C to stabilise the microstructure and minimise the influence of precipitation hardening. 2.2. Microstructural characterisation

n

Corresponding author. Tel.: þ 43 1 5880130867; fax: þ43 1 58801 30899. E-mail address: [email protected] (R. FernándezGutiérrez). 0921-5093/$ - see front matter & 2014 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.12.093

The microstructure of the alloy was analysed by light optical microscopy and scanning electron microscopy before and after creep exposure.

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The α-Al was leached out using an 18.5 vol% solution of HCl during 15 min to reveal the morphology and interconnectivity of the eutectic Si. 2.3. Creep tests Cylindrical samples with a gauge length of 36 mm and a diameter of 6 mm in AC, 20 min ST and 64 h ST conditions were subjected to isothermal tensile creep tests under constant load conditions at 300 1C in a range from 20 to 50 MPa. Creep results of the 4 h ST condition are taken from previous investigations [17]. The strain was measured using linear variable displacement transducers with a resolution of 0.01 mm and recorded by a digital acquisition system.

3. Results

Fig. 2(a)–(d) shows scanning electron micrographs of the eutectic Si network after leaching out the α-Al matrix from the alloy in AC, 20 min, 4 h and 64 h ST conditions, respectively. The lamellar multilayer three dimensional Si structure observed in AC condition becomes rounder and more fibrous-like with increasing ST time. The interconnectivity seems to be at least partially conserved after ST. The evolution of the volume fraction of intermetallic phases during ST was quantified from three different optical micrographs of the microstructure in each condition. Fig. 3 typifies the process applied to separate them from the rest of the microstructure. From the characteristic light optical micrograph, Fig. 3(a), the intermetallic phases were manually coloured, Fig. 3(b). Finally a binarised image (black and white) was produced using the ImageJ software [22], Fig. 3(c). The quantitative results for all the conditions can be observed in Fig. 4. The volume fraction, Vf, of intermetallics decreases sharply during the first 20 min of ST and dissolution continues in a lower proportion even up to 64 h ST at 480 1C.

3.1. Microstructure Fig. 1(a)–(d) shows light optical micrographs of the AlSi12CuMgNi alloy in AC condition after 20 min, 4 h and 64 h ST at 480 1C, respectively. The following intermetallic phases were identified based on energy-dispersive X-ray analysis and correlation with literature [20]: (Mn–Fe)3Si2Al15, FeNiAl9, Al2Cu, Mg2Si, Al7Cu4Ni (see Fig. 1a). An aluminide containing Mg, Ni, Fe and Si was also found which corresponds to a modification of the compound Al8FeMg3Si6 [21]. This aluminide and Mg2Si dissolve during solution treatment. A clear spheroidisation of the eutectic Si is observed with increasing ST time. Table 1 Alloy composition (wt%). Element

Al

Si

Cu

Mg

Ni

Fe

Mn

Zn

Ti

Mass [%]

Matrix

11–13

0.8–1.3

0.8–1.3

1.3

0.7

0.3

0.3

0.2

3.2. Creep tests Representative strain vs. time, strain rate vs. time and strain rate vs. strain curves obtained during the creep tests carried out at 40 MPa are shown in Fig. 5(a)–(c), respectively. The curves, with some experimental scattering, show a typical primary creep behaviour at the beginning of the experiments with the strain rate decreasing until a minimum creep rate is reached. Afterwards, the strain rate steadily accelerates with increasing creep strain (Fig. 5c) until sample failure. The effect of ST on creep resistance can be clearly observed in the figures. The minimum creep rate, ε_ min , of each experiment is plotted in Fig. 6 as a function of the corresponding load in a doublelogarithmic scale. The AC condition shows lower minimum creep rates and longer creep lifetime than those of the ST conditions for all the applied loads. The 4 h and 64 h ST conditions show similar ε_ min within the experimental scattering. The minimum creep rate

Fig. 1. Light optical micrographs of the AlSi12CuMgNi alloy in (a) as cast condition and after solution treatment at 480 1C for (b) 20 min, (c) 4 h and (d) 64 h.

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Fig. 2. Scanning electron micrographs of the deep etched AlSi12CuMgNi alloy revealing the silicon network: (a) as cast condition, after solution treatment at 480 1C for (b) 20 min, (c) 4 h and (d) 64 h.

Fig. 3. Process of segmentation of intermetallic phases from light optical micrographs. This example shows (a) an original micrograph of the AlSi12CuMgNi alloy in as cast condition, (b) the intermetallic phases hand-painted in pink and (c) the binarised image with the intermetallic phases in white and the rest of the microstructure in black. (For interpretation of references to colour in this figure, the reader is referred to the web version of this article.)

Fig. 4. Evolution of the volume fraction of intermetallic phases during solution heat treatment at 480 1C.

for the AC and 20 min ST conditions and decreases to  4.5 after 4 h and 64 h of solution treatment. This evaluation differs from that presented in [17] for the 4 h ST condition but it is considered more accurate based on the results obtained in the present work for the other investigated conditions. Fig. 7(a) and (b) shows light optical micrographs of samples in AC condition tested until failure at 30 and 50 MPa, respectively. The load direction is the horizontal axis. Damage features characteristic for all tested conditions can be observed. Fig. 7(a) shows a secondary crack located at  500 mm from the fracture surface. The crack developed perpendicularly to the load direction through a region of eutectic Si revealing the brittle nature of the damage process. Fig. 7(b) shows a detail of a region close to the fracture surface where microfracture of the eutectic Si and the aluminides can be observed. A few sites where debonding between the Almatrix and the rigid phases took place are indicated. Void formation was not observed in the Al-matrix. These damage types can be found for all investigated conditions, although fracture of aluminides is more evident at higher applied loads.

4. Discussion at 20 MPa for the 20 min ST condition lies between the values obtained for AC and 64 h ST. For loads Z30 MPa the minimum creep rate for 20 min ST coincides with 64 h ST. The creep exponent, n, was determined for each condition applying a linear fit to the experimental ε_ min values in Fig. 6. The results are shown next to the symbol labels in this figure together with the standard error and coefficient of determination of the linear fit, R2. n is 6.6

4.1. Effect of solution treatment on the creep behaviour The creep results shown in Figs. 5 and 6 reveal that the AlSi12CuMgNi alloy in as cast condition presents the highest creep resistance among all studied conditions. The minimum creep rate gradually increases with solution treatment time. On the other

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Fig. 5. Creep test results obtained for an applied load of 40 MPa: (a) strain vs. time, (b) strain rate vs. time and (c) strain rate vs. strain.

Fig. 6. Dependence of ε_ min on s for the AlSi12CuMgNi alloy in as cast condition and after exposure at 480 1C for 20 min, 4 h and 64 h.

hand, the creep exponent n decreases from 6.6 to 4.4 between the as cast condition and 4 h ST and remains constant between 4 and 64 h of solution treatment at 480 1C. The differences in the creep behaviour between the investigated conditions can be attributed to the microstructural changes taking place during the solution treatment: dissolution of intermetallic compounds and spheroidisation of the eutectic Si. The latter comprises of the disintegration of three dimensional Si network and further rounding of disintegrated particles [23]. The present investigations revealed the rounding of the eutectic Si (see Fig. 1). Three dimensional studies of the microstructure of similar piston alloys and cast AlMgSi alloys indicate that the largest decrease in the interconnectivity of the eutectic Si takes place during the first 1–4 h of solution treatment [13,24–26]. Table 2 summarises creep data for similar cast Al–Si alloys [14–18]. The experimental conditions utilised for the creep experiments (alloy,

heat treatment, load mode, test temperature and load range) are presented together with the obtained creep exponents n. The creep exponents for TZ300 1C are between  4.4 and 6.6 for experiments carried out under tensile load conditions, which are in full agreement with the results obtained in the present work. A close look to the heat treatments in Table 2 reveals the reason for the different reported creep exponents. Jeong et al. [18] studied alloys subjected to an ageing heat treatment at 250 1C without any previous solution heat treatment. This is analogous to the AC condition in the present work. The AlSi12Ni1.2Cu alloy is the closest in composition to AlSi12CuMgNi among the materials studied in [18]. The creep exponent for this alloy is the same as the one obtained in the present work for the AC condition (n¼6.6). On the other hand, the data reported by Bidlingmaier et al. [15] for an AlSi12CuMgNi alloy subjected to 1 h ST at 480 1C plus overaging show the same creep exponent n¼4.4 as after 4 h and 64 h ST in the present investigations. Unfortunately, Dlouhy et al. [14] did not report the solution treatment time used during their investigations. The creep exponent obtained in [14] for AlSi7Cu3 and AlSi7Cu3Mg is n 6, which may be an indication of solution treatment time o1 h. The different creep exponents reported by Cseh et al. [16] with respect to the other investigations may be due to the different type of load applied for the creep tests (indentation creep). It can therefore be concluded that the change observed for the creep exponent n is a function of the solution heat treatment applied to the AlSi12CuMgNi and the microstructural changes taking place during the first hour of ST. 4.2. Load transfer loss during solution treatment An analogy with MMCs can be done for the investigated alloy since the eutectic Si and the intermetallic compounds do not creep at 300 1C. Therefore, each ST condition can be considered as an individual MMC with different reinforcement volume fractions and reinforcement morphologies. These rigid phases are able to carry part of the applied load, thus decreasing the stress seen by the α-Al matrix during creep exposure. The creep behaviour of

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Fig. 7. Characteristic creep damage features observed for the AlSi12CuMgNi alloy: (a) secondary crack through eutectic Si (  500 mm from the fracture surface) in an as cast sample tested at 30 MPa, and (b) region close to the fracture surface showing microfracture of the eutectic Si and the intermetallic phases as well as debonding between the matrix and the rigid phases (as cast condition tested at 50 MPa).

Table 2 Summary of creep studies of cast Al–Si alloys. Author, year

Material

ST T/tþ Ageing T/t

Heat Load mode Test treatment temp. (1C)

Load range Creep (MPa) exponent n

Remarks

Dlouhy et al. (1993) [14]

AlSi7Cu3

500 1C/n.a. h þ 160 1C/n.a. h

T6

10–55

Comparison with short-fibre reinforced MMC

Tensile

350

AlSi7Cu3Mg

6.3

Bidlingmaier et al. AlSi12CuMgNi (1996) [15]

480 1C/1 h þRT/ 24 h þ 215 1C/9 h

T7

Tensile

Cseh et al. (1999) [16]

AlSi12CuMgNi

480 1C/10 h (oil q.) þ RT/ 100 h þ350 1C/24 h

T7

Requena et al. (2006) [17]

AlSi12CuMgNi

480 1C/4 h  (oil q.) þ 190 1C/ T7 4 h þ 300 1C/3 h

Jeong et al. (2010) [18]

AlSi12Ni1.2Cu 250 1C/5 h AlSi12Ni2.63Cu2.85 AlSi12Ni2.83Cu4.89

T5

MMCs, in which an effective load transfer from the matrix to the reinforcement takes place, can be described by the following relationship [27].

ε_ min ¼ Aðs  sT Þn ¼ AΔsn

6.0

ð1Þ

where A is a constant, sT is the load transferred from the matrix to the reinforcement at the applied load s, and Δs is the creep strength increment [27,28]. In the present case, the lowest creep resistance is shown by the 64 h ST condition. Although this condition may still present some reinforcement effect by the eutectic Si and the intermetallic compounds, it will be taken as a reference to evaluate the relative load transfer loss after solution treatment. This is opposite to the creep strength increment defined in [27]. The experimental scattering shown by the creep results (see Figs. 5 and 6) make it difficult to estimate the load transfer loss for the entire load range but an indicative trend can be estimated between 30 and 40 MPa. For this, lines with slopes n þ Δn and n  Δn were plotted intersecting the linear regressions of AC, 20 min and 64 h (Fig. 6) at 20 MPa and 50 MPa. Then, the maximum and minimum load transfer loss of the 64 h ST condition with respect to AC and 20 min ST were calculated as the maximum and minimum possible load differences between these lines. The load transfer loss at 30, 35 and 40 MPa after 64 h of solution treatment was then

350

n.a.

4.4

Reference matrix for comparison with short-fibre reinforced MMC

Indentation 249 276 301 324 348

30–130

7.63 7.41 6.51 6.52 6.5

Tensile

300

20–40

3 between 20–35 MPa

n¼ 4.5 70.5 if load range 20–40 MPa is considered

Tensile

400

10–30

6.6 5.9 5.3

Addition of Cu and Ni reduces the strain rate

estimated as the mean of these previously calculated load differences and the one obtained between the linear regressions (lines with slope ¼n). These values did not show significant discrepancies between 30 and 40 MPa. Therefore, the average load transfer loss of the 64 h ST condition with respect to AC and 20 min is shown in Fig. 8. The load transfer loss between the AC and 64 h ST conditions is 8.6 MPa70.8. On the other hand, the load transfer loss amounts 3.2 71 MPa between 20 min and 64 h ST. The difference between these values (  5.4) is the load transfer loss in the first 20 min of solution treatment. This indicates, and supports the discussion in the previous section, that the first 20 min to 1 h of solution treatment result in the largest decrease of the creep resistance of the AlSi12CuMgNi alloy. Following the analogy with MMCs, it is possible to consider the investigated alloy as an α-Al matrix reinforced with an equivalent volume fraction of short fibre reinforcement with a certain aspect ratio. This can be done applying the solution proposed by Fernández et al. [27] using the shear-lag model modified for short fibres with an effective aspect ratio, Seff, as shown by Ryu et al. [29], Seff ¼

2½ð1  Vf r Þs  sef f  Vf r sef f

ð2Þ

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drastically its load bearing capability [25] and, consequently, the creep resistance of the alloy.

5. Conclusions The creep resistance of an overaged AlSi12CuMgNi alloy produced by squeeze casting was studied as a function of solution treatment time at 480 1C by means of isothermal creep tests at 300 1C. The following conclusions can be drawn from the analysis of the results:

Fig. 8. Load transfer loss of the 64 h ST condition with respect to the AC and 20 min ST conditions.

Table 3 Effective aspect ratio Seff of equivalent short fibre reinforced composites with respect to the 64 h and 20 min ST conditions. AC

s ¼ 30 MPa 19.9 s ¼ 35 MPa 13.9 s ¼ 40 MPa 9.6 Reference¼64 h ST

20 min ST

AC

8.6 6.9 5.7

21.6 17.9 15.0

Reference¼ 64 h ST

Reference¼ 20 min ST

where s is the load applied to the composite, seff is the effective load carried by the matrix and Vfr is the volume fraction of reinforcement. This equation only considers the effect of the Vfr with respect to an unreinforced matrix. Eq. (2) can be modified to obtain the effective aspect ratio of the reinforcement in a composite with respect to another reference composite with a lower reinforcement volume fraction, Seff ¼

2½ð1  ΔVf r Þs  sef f  ΔV f r sef f

ð3Þ

ΔVfr is the difference in reinforcement volume fraction which in the present case is the difference in the volume fraction of intermetallic phases between the conditions considered since the volume fraction of Si remains constant during solution treatment. This means that the reinforcement volume fractions for AC and 20 min ST conditions with respect to 64 h are ΔVfr  4 vol% and 2.2 vol%, respectively (see Fig. 4). Similarly, the difference in reinforcement volume fraction between the AC and 20 min ST condition is ΔVfr  1.8 vol%. Furthermore, Eq. (3) implicitly contains in s the effect caused by the change in morphology and interconnectivity of the eutectic Si since these are the only further changes taking place in the microstructure of the alloy. The absolute values obtained for Seff using Eq. (2) do not mean that the studied alloy is a short fibre reinforced composite with the respective aspect ratio in each condition. This factor only shows the aspect ratio that equivalent short fibre reinforced metal matrix composites should have to show the same difference in creep resistance as observed experimentally in this work. Table 3 shows the effective aspect ratio Seff for different applied loads with respect to 64 h ST and 20 min ST conditions. Although the change in reinforcement volume fraction is  2 vol% between the AC – 20 min and between 20 min and 64 h ST conditions, the Seff is  3 times larger for the AC – 20 min ST case. Again, this reflects that the first part of the solution treatment (o1 h) affects mostly the loss of interconnectivity of eutectic Si reducing

1. The solution treatment results in the partial dissolution of intermetallic compounds from  10.4 vol% in AC condition to  6.4 vol% after 64 h ST. Also, spheroidisation of eutectic Si was observed in the same period. 2. The highest creep resistance is shown by the AC condition and gradually decreases with increasing solution treatment time. This is attributed to the loss of load bearing capability of the rigid phases (intermetallic compounds and eutectic Si). This is due to the microstructural changes taking place during solution treatment, namely the partial dissolution of intermetallic compounds and the spheroidisation of eutectic Si. 3. The determination of the load transfer loss during solution heat treatment and an analogy with a modified shear-lag model for metal matrix composites indicate that the largest reduction in creep resistance occurs in the first 20 min–1 h of solution heat treatment. Comparison with results from the literature for similar alloys indicates that this may be driven by the loss of interconnectivity of the eutectic Si network.

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