Quench Sensitivity of an Al-7%Si-0.6%Mg Alloy: Characterization and Modeling

Transcription

1 Quench Sensitivity of an Al-7%Si-0.6%Mg Alloy: Characterization and Modeling Murat Tiryakioğlu 1 Robert Morris University Department of Engineering Moon Township, PA USA Ralph T. Shuey Alcoa Technical Center Alcoa Center, PA USA Keywords: Mg 2 Si, C-Curves, Quench Factor Analysis, Heat Treatment, Modeling Abstract The quench sensitivity of a cast Al-7wt.%Si-0.6wt.%Mg alloy was characterized by tensile tests and scanning electron microscopy. Specimens were cooled from the solution treatment temperature following 58 different cooling paths including interrupted and delayed quenches. Analysis of microstructure showed that quench precipitates were Mg 2 Si (β) which nucleated heterogeneously on Si eutectic particles as well as in the aluminum matrix, presumably on dislocations. The quench sensitivity of the alloy s yield strength was modeled by multiple-c curves, using an improved methodology for quench factor analysis. The three C-curves used in the model represented loss of solute by (i) diffusion of Si to eutectic particles, (ii) precipitation of β on Si eutectic particles, and (iii) precipitation of β in the matrix. The model yielded a R 2 of and a Root Mean Square for Error (RMSE) of 7.4 MPa. The model and the implications of the results are discussed in the paper. 1 Corresponding author Tiryakioglu@rmu.edu 1

2 Introduction Cast Al-7wt.%Si-Mg alloys are in-situ metal matrix composites (MMC), due to the presence of eutectic Si particles. They have been used successfully in a variety of automotive and aerospace applications. Heat treatment of these alloys involves a solution treatment, subsequent quenching, and finally natural and/or artificial aging. During solution treatment, the Mg 2 Si (β) phase is dissolved, the Al matrix is homogenized and the eutectic (Si) phase becomes dissociated and spheroidized. During artificial aging, β precipitates form, which subsequently increase the yield strength. Quenching the castings at the highest rate possible, as in cold water, retains more solute and vacancies in solution, which increases the yield strength attainable after aging. However, high cooling rates associated with water quenching result in the generation of thermal stresses leading to distortion and residual stresses, especially in castings with complex geometries. Conversely, low cooling rates that provide reduced levels of thermal stress issues, produce non-strengthening quench precipitates, such as β and β in the Al-Mg-Si system, which ultimately reduce the strength attainable after aging. Therefore process engineers strive to design cooling processes with an optimum balance of strength and thermal stresses. To accomplish this task, the effect of different cooling paths on the strength needs to be characterized and modeled. Furthermore, the dependences on composition, homogenization and aging need to be understood and ultimately modeled. To attain this manifold, quantitative, mastery of process-property relations, investigation should follow the intermediate microstructure, i.e., the precipitates formed during quench. Several studies on the quench sensitivity of cast Al-7%Si-Mg alloys have been reported [1,2,3,4,5,6,7]. In only three of these [1,2,3] was microstructure investigated for quench precipitates. The literature on quench sensitivity of wrought Al-Mg-Si alloys, is relatively voluminous, but use for Al-7%Si-Mg casting alloys requires identifying and understanding the differences. The following advances are now reported: quench sensitivity of D357, an aerospace casting alloy, was investigated using a variety of interrupted quench and delayed quench paths and metallographic examination. Experimental findings are compared with the limited prior work on Al-7%Si-Mg alloys. A quench model is presented, which generalizes previous versions to include multiple types of quench precipitates, and diffusion to eutectic Si during quench. The quench sensitivity of D357 is then compared with that of two wrought Al-Mg-Si alloys. 2

3 Previous Work Quench Sensitivity of Yield Strength Tsukuda et al. [7] investigated the effect of quench water temperature on the yield strength of A356 (0.30 wt.% Mg) alloy specimens and found that water temperatures between 15 and 90 o C did not affect σ Y. Tsukuda et al. also found that a quench delay of 120s in still air (delayed quench at 460 o C) produced a decrease of 14 MPa from the peak strength of 169 MPa obtained after artificial aging at 155 o C for 4.5 hours. These results are in contrast with those of Zhang and Zheng [1] who found a sharp decrease in σ Y with decreasing cooling rate. Zhang and Zheng attributed this contrast to the difference in aging treatments between the two studies; the alloy was underaged in the study by Tsukuda et al., whereas it was peak-aged in the study by Zhang and Zheng. The underaged condition was less quench-sensitive than the peak-aged condition. The effect of aging treatment on the quench sensitivity of A356 (0.40 wt.% Mg) was examined by Croucher and Butler [6] who quenched specimens in water at three different temperatures and aqueous polyalkylene glycol solution with 5 different concentrations. Specimens were then artificially aged at 154 o C for 10 hours or at 177 o C for 3 hours. Results showed that the decrease in σ Y with cooling rate between o C, and hence quench sensitivity, is less for specimens aged at 154 o C for 10 hours. Rometsch and Schaffer [5] studied the quench sensitivity of A356 and A357 alloys (with 0.40 and 0.62wt.%Mg, respectively) by conducting continuous cooling experiments in 7 (5 for A357) different cooling media. Specimens were solution treated at 540 o C and artificially aged immediately after quenching for 8 hours at 170 o C. Rometsch and Schaffer used Quench Factor Analysis (QFA) to model the quench sensitivity of both alloys. QFA was developed previously by Staley and coworkers [8,9,10] as a process-property model. Results of QFA predicted the 90% iso-σ Y curves for A356 and A357 to have noses at 6 and 15 seconds, respectively, at 350 o C. These results indicate that A356, despite its lower solute content, has higher quench sensitivity than A357, which is in contrast with the principles of physical metallurgy. The reason for this anomaly is speculated to be due to low number of cooling paths used in the study. Despite much scatter in hardness results, Robinson [2] plotted iso-vhn plots for the three alloys; the alloy with 0.59% Mg had a 90% iso-vhn curve with a nose at approximately 350 o C and 10 seconds. The location of the nose is similar to the one reported by Rometsch and Schaffer for A357. 3

4 Zhang and Zheng [1] reported that the quench sensitivity is higher in the cast alloy than in a wrought alloy with similar Mg 2 Si content. This increased quench sensitivity was speculated to be due to (i) due to excess Si in solution and (ii) the presence of Si eutectic particles, which may act as nucleation sites. Quench Precipitates Zhang and Zheng [1] investigated the quench sensitivity of an Al-7%Si-0.4%Mg (A356) alloy by quenching specimens in 4 different media. They calculated average quench rates by dividing the temperature fall by the time elapsed when the specimens cool down from 450 to 200 o C. The average quench rates ranged from 250 to 0.5 o C/s. Specimens were solution treated at 540 o C for 14 hours and artificially aged for 6 hours at 170 o C without any aging delay. Analysis of microstructure via transmission electron microscopy revealed β precipitates in the Al-matrix of slowly-cooled specimens, as presented in Figure 1. Note that β precipitated along dislocations and that a precipitate free zone (PFZ) formed around the β precipitates. Robinson [2] conducted interrupted quenching experiments with Al-7%Si alloys with three Mg levels (0.26, 0.38 and 0.59 wt.%). Specimens were solution treated for 14 hours at 540 o C and artificially aged for 6 hours at 170 o C within 15 minutes of cooling to room temperature. Robinson analyzed the microstructure via transmission electron microscopy to find the reason for hardness drop during isothermal holding at 450 o C. Transmission electron micrographs showed no quench precipitates in the aluminum matrix. Consequently, Robinson speculated that the reason for drop in hardness is the loss of Si in solution by its diffusion to eutectic Si particles. In addition to the loss of Mg 2 Si in the form of quench precipitates, the loss of excess Si is also an important contributor to the quench sensitivity of cast Al-7%Si-Mg alloys; it is established that the excess Si in Al-Mg-Si alloys promotes an additional response to age hardening by both refining the size of the strengthening (β ) particles [11] and precipitating as Si [12]. In an Al-7%Si-0.2%Mg alloy, Pedersen and Arnberg [3] observed coarse Si quench precipitates within Al dendrites and PFZs near eutectic regions, as shown in Figure 2. The PFZ around the eutectic particles was attributed to the loss of Si by diffusion to the eutectic particles. In another study, Pedersen and Arnberg [13] investigated the microsegregation of Si within Aldendrites by microprobe analysis and found that peak concentration was in the midpoint of the dendrites. These results were later verified by simulations of Dons et al. [14], who concluded 4

5 that the anomalous microsegregation pattern was developed after solidification, only in specimens that have solidified at a high rate. In addition, cooling rate to room temperature has to be low. This explains the results of Zhang [15] and Chen et al. [4]. Zhang investigated the precipitation of excess Si in A356 (0.38%Mg) after a solution treatment of 14 hours 540 o C, followed by water quench or cooling in still air, and artificial aging at 170 o C for 6 hours. In specimens quenched in water, Si precipitated during aging, mainly in the eutectic region and occasionally in the Al-matrix on dislocations. In specimens cooled in still air, however, no Si precipitates were observed. Zhang attributed the absence of Si precipitates to the diffusion of Si in solid solution to eutectic Si particles. Chen et al. [4] observed that the morphology of the Si phase in A357 alloy (Al-7%Si-0.6%Mg) changed with decreasing quench rates, implying the growth of Si particles during cooling from solution treatment, which is consistent with the results of other studies. The literature on the quench sensitivity of cast Al-Si-Mg alloys indicates that there is loss of Si in solution by diffusion to Si particles at 450 o C, there are no quench precipitates in the aluminum matrix β precipitates along dislocations in the matrix at low quench rates in A356 aging practice affects quench sensitivity cast Al-7%Si-Mg alloys may be more quench sensitive than wrought alloys with the same solute content. The literature survey above has also shown that the following issues remain unattended: possible quench precipitates on Si particles the nature of matrix precipitates in D357 the characterization, via a process-structure-property model, of how different quench precipitates affect yield strength obtained after aging a systematic comparison of quench sensitivity of cast Al-7%Si-Mg and wrought 6XXX alloys with similar solute contents. 5

6 Experimental Details D357 alloy (not modified by Na or Sr) tensile specimens excised from plates cast by Hitchcock Industries, Inc., Minneapolis, MN, were used in our experiments. The chemical composition of the alloy is given in Table 1. The details of the filling and feeding system used in the mold are unknown. The geometry of tensile specimens is given elsewhere [16]. The specimens were solution treated in an air-recirculating furnace at 540 o C for 8 hours. Solution treatment was followed by 58 different quench paths to obtain a wide interval of yield strengths. Quenches were interrupted at 4 temperatures (450, 400, 350 and 300 o C) for durations ranging from 3 to 100,000 seconds. In delayed quench experiments, specimens were initially cooled in four different media (boiling water, 20% aqueous polyalkylene solution, forced air and still air) until the target temperatures were reached, and subsequently quenched in cold water. Cooling curves for the delayed quench experiments are presented in Figure 3. Time-temperature data were collected by a data acquisition system connected to type-k thermocouples placed in the geometric center of dummy tensile specimens. In addition, one specimen was quenched in cold water, and 8 specimens were cooled in furnace to different temperatures followed by cold water quenching. Specimens were then left to naturally age at room temperature for 24 hours before they were artificially aged at 170 o C for 8 hours. An Instron tensile tester was used at an engineering strain rate of 0.001/s and 0.2% offset yield strength (σ Y ) values were recorded. Certain interrupted quench specimens were investigated metallographically using a scanning electron microscope (SEM) equipped with a field emission gun (FEG) at Alcoa Technical Center. 6

7 Results The effect of isothermal holding time on yield strength at various temperatures is presented in Figure 4. Note that the highest rate of loss in yield strength takes place at 350 o C. Figure 4 also indicates the specimens chosen for microstructural investigation. Scanning electron micrographs of the selected specimens are shown in Figures 5-8. All specimens included a small number of large and generally isolated particles interpreted as undissolved after solution treatments. Thermodynamic calculation gave solvus temperature for the alloy composition as 547 ºC, so incomplete solutionization is understood in hindsight. Also present in all specimens were more numerous, smaller particles interpreted as quench precipitates nucleated on eutectic Si. No direct evidence on the diffusion of Si in solution to eutectic particles during quench was found. With high magnification (Figures 7.b and 8.b), fine quench precipitates in the matrix at lower hold temperatures were identified. The size and distribution are consistent with the β precipitates reported by Zhang and Zheng [1]. However, they do not show the characteristic crystallographic orientation of β, which might be attributed to poor image quality or possibly to conversion from β to β during the isothermal hold. Modeling Quench Sensitivity The Model The mathematical model used is an improved version of the Quench Factor Analysis (QFA) approach. The improved model establishes process-structure-property relationships and accommodates multiple quench precipitates and uses thermodynamic and calculated data for certain coefficients. The amount of each quench precipitate is represented by a unitless microstructural state variable S. For arbitrarily long isothermal hold, in the absence of competing precipitates, the limiting value, S eq, is S eq H 1 1 exp R K 4 T (1) T 1 ; T K 4 where K 4 is the solvus temperature and ΔH the precipitation enthalpy for quench precipitate. S eq represents the equilibrium amount of S at a given temperature. The evolution of quench precipitates is modeled by 7

8 ds S dt where t c is critical time to nucleate and grow quench precipitates, and is given by [8-10] eq t S c (2) 2 K 3K4 K5 tc K2 exp (3) 2 RT K T RT 4 where K 2 : coefficient related inversely to number of nucleation sites (s) K 3 : coefficient related to energy barrier to heterogeneous nucleation (J/mol) K 5 : stochiometrically-averaged activation energy for diffusion (J/mol) R: gas constant, J/(mol-ºK) T: temperature (ºK) The numerical algorithm is S At the end of the quench, S is found by i t i ( S eq Si 1) 1 exp (4) t c S S (5) i i Yield strength can then be estimated by S (6) Y max k j j j where j is for the quench precipitates modeled, σ max is the strength obtained at an infinite quench rate and k j is the strength coefficient. Equation 6 assumes that yield strength depends linearly on the amount of solute available after quench, which is consistent with past studies [17,18,19]. The model presented in this study is an expansion of previous models [8,9,10,19,20] 8

9 which considered only one type of quench precipitate (same stochiometry and site) by using similar equations. In generalizing to multiple quench precipitates, the intention was to maintain compatibility with those previous models. With a single quench precipitate, Equations 1-6 are equivalent to those given in previous studies. The Application of the Model to D357 Data For each quench precipitate, K 2, K 3 and k were varied along with σ max. Hence a total of 10 variables were fitted to data using Microsoft Excel Solver. The coefficients fitted are reported in Table 2. K 4, K 5 and ΔH were calculated from thermodynamics and previously published data. Note that K 3 for Si is 0, making the curve effectively pure diffusion or alternatively, pure growth of eutectic Si particles. Actually the shape of a C-curve for pure growth differs from that for nucleation and growth with zero energy barrier to nucleation (i.e., K 3 = 0), but such refinement seems unwarranted. Critical times (t c ) for the three quench precipitates versus temperature are shown in Figure 9, which also shows the hold times and temperatures that produced the microstructures in Figures 5-8. Experimental versus predicted σ Y results are presented in Figure 10, showing excellent agreement between the two. The root mean square error (RMSE) for the fit was 7.4 MPa. An alternative model with only the two C-curves for β precipitation yielded a RMSE of 9.9 MPa. The triple-c fit is preferred not only for its lower RMSE, but also because of the strong physical evidence on Si loss to eutectic particles provided in the literature. In addition, based on the findings of Dons et al. [14] diffusion of Si in Al dendrites is expected to take place for slow cooling rates since solidification rate was probably quite high due to presence chills in the mold. Figure 11 shows predicted isothermal decay curves at the four temperatures. The curves are again in excellent agreement with experimental data. An interesting feature of these curves is that at long holding times, there seems to be a second drop, due to loss of Si to existing Si particles. Iso-σ Y contours are presented in Figure 12. 9

10 Discussion The model above, like those using original QFA methodology [8-10], enables prediction of strength loss for arbitrary quench path. In addition there are two novel aspects; (i) multiple precipitates are represented by separate C-curves, and (ii) K 3 =0 for loss of Si in solution by diffusion to existing Si particles. The multiple C-curves shown in Figure 9 also explain why Robinson failed to find any quench precipitates within Al dendrites; at 450 o C, β precipitates only on Si particles and not in the matrix. A C-curve for precipitation of Si in the aluminum matrix at lower temperatures can be expected. However no Si precipitates were observed in our microstructural analyses. If Si precipitates are present in the samples, they are expected to have precipitated during aging and therefore to be fine and strengthening. As indicated previously, some undissolved β-phase was observed in all specimens. This is due to the solution treatment temperature (540 o C) being less than the solvus temperature (547 o C) for the alloy composition, leaving 0.04% Mg and 0.07%Si undissolved. If the specimens were solution-treated at or above 547 o C, K 4 would be 820 K and σ max would increase slightly with the other coefficients in Table 2 remaining the same. The model presented in this study applies to the composition with Mg and Si reduced by their respective undissolved amounts, then completely solutionized at 540 o C or higher. The iso-σ Y contours presented in Figure 12 shows that the nose of the 90% iso-σ Y curve is at 345 o C and at 25 seconds. The temperature of the nose is comparable to that reported by Rometsch and Schaffer [5] and Robinson [2]. However, the nose of the curve is at a longer time than the ones reported Rometsch and Schaffer (15 s) and Robinson (10 s). Hence, the quench sensitivity of D357 is less than the two alloys with similar Mg contents in the two previous studies. This might be attributed to the low number of cooling paths used by Rometsch and Schaffer and high variability in VHN in the results of Robinson. Another possible factor is the two-step aging practice with natural aging as first step, employed in this study, and not in the other two studies. A reduction of quench sensitivity with increasing aging delay (natural aging) before artificial aging was reported by Taylor [21] on 6351 (0.60 wt.%mg, 0.99wt.%Si) at 170 o C. The peak yield strength after artificial aging decreased slightly with increasing aging delay (up to 3 days) for specimens quenched at a high rate (>1000 o C/s). Aging delay had a beneficial effect in specimens quenched very slowly (0.65 o C/s). Together these trends imply that peak 10

11 strength with natural aging is less sensitive to quench than is peak strength with no delay before artificial aging. Comparison of Quench Sensitivity of Cast Al-7%Si-Mg and Wrought 6XXX Alloys The 95%-σ Y contours for D357 and two wrought alloys are presented in Figure 13. Contours for the two wrought alloys were drawn from the raw data of Bergsma et al. [22] for 6061-T6 (0.65wt.% Si, 0.89wt.%Mg, 0.23wt.%Cu) and Zajac et al. [23] for 6082-T6 (0.92wt.%Si- 0.59wt.%Mg). Figure 13 indicates that 6061, which has much more solute is only slightly more quench sensitive than D , however, with almost the same Mg 2 Si content as D357, is much less quench sensitive than D357. These results are in agreement with the observations of Zhang and Zheng. This high quench sensitivity of cast Al-Si-Mg alloys is mainly due to the presence of Si particles in the structure, which has affects quench sensitivity in three ways: (i) Si in solid solution diffuses to these particles, (ii) particles serve as heterogeneous nucleation sites for β, and (iii) dislocation density in dendrites of Al-7%Si alloys are expected to be high due to the mismatch of coefficient of thermal expansion (CTE) between Si and Al. Large differences in CTE between Al-matrix and Si particles result in generation of compressive stresses in the particles and tensile stresses in the matrix during cooling from high temperatures. Some of the CTE mismatch is accommodated by plastic relaxation in the matrix, increasing the dislocation density of the matrix [24], which provides more nucleation sites for quench precipitates (decreased K 2 in Equation 3). This assumption was verified in wrought alloys reinforced with ceramic particles. Dutta et al. [25] studied the effect of Al 2 O 3 additions as reinforcement particles on the aging response of The reinforced specimens had a significantly higher dislocation density in the matrix, which accelerated both the nucleation and growth of precipitates. That explains why Papazian [26] found that SiC additions increased the quench sensitivity of Similar results were reported for reinforced 2xxx-series alloys [27,28]. 11

12 Conclusions The quench sensitivity of D357 alloy is due to (i) the loss of Si to eutectic particles, (ii) precipitation of β on eutectic Si particles, and (iii) precipitation of β in the aluminum matrix. Natural aging before artificial aging (aging delay) may decrease the quench sensitivity of D357. D357 is slightly less quench sensitive than 6061, which has much higher solute, and more quench sensitive than 6082 with almost the same Mg 2 Si content. The high quench sensitivity of D357 is due to the presence of Si particles. Multiple C-curves need to be employed to effectively model the quench sensitivity of cast Al-Si-Mg alloys. Acknowledgements MT acknowledges a Summer Research Fellowship from Robert Morris University. 12

13 Table 1. The chemical composition (in wt.%) of D357 used in this study. Si Mg Cu Fe Be Ti Al Rem. Table 2. Coefficients for the D357 quench model. Shaded coefficients were determined from thermodynamics and existing diffusion data, others by fitting the model to experimental data. K 2 K 3 K 4 K 5 k ΔH (s) (J/mol) (K) (J/mol) (MPa) (J/mol) (MPa) Si Diffusion to Particles , , β on Si Particles , ,066 β in the Matrix , ,066 σ max 13

14 Figure 1. β quench precipitates in a slowly-cooled A356 alloy [1] (Courtesy of Metallurgical and Materials Transactions A). 14

15 Figure 2. Precipitate free zones near eutectic Si particles in a slowly-cooled Al-7%Si alloy [3]. The PFZ was attributed to the loss of dissolved Si by its diffusion to Si particles (Courtesy of Metallurgical and Materials Transactions A). 15

16 T ( o C) T ( o C) M. Tiryakioǧlu, R.T. Shuey: Metall. Mater. Trans. B, v. 38, pp , time (s) (a) time (s) (b) Figure 3. Cooling curves in delayed quenching of D357 in various media: (a) boiling water and 20% aqueous polyalkylene glycol solution (indicated by arrows), (b) forced and still air (indicated by arrows). 16

17 C 400 C 350 C 300 C 250 Y (MPa) ,000 10, ,000 1,000,000 Hold Time (s) Figure 4. The effect of hold time at different temperatures during interrupted quenching on the yield strength of D357. Specimens selected for microstructural investigation are indicated by arrows. 17

18 Si β Figure 5. Microstructure of the specimen held for 10,000 seconds at 450 o C. 18

19 (a) Figure 6. Microstructure obtained after holding specimens at 400 o C for (a) 30 s and (b) 200 s. (b) 19

20 (a) (b) Figure 7. Microstructure of the specimen held at 350 o C for 100s showing β precipitates (a) on Si particles, and (b) in the aluminum matrix. 20

21 (a) Figure 8. Microstructure of the specimen held at 300 o C for 100s β precipitates (a) on Si particles, and (b) in the aluminum matrix. (b) 21

22 T ( o C) M. Tiryakioǧlu, R.T. Shuey: Metall. Mater. Trans. B, v. 38, pp , ST Temperature Diffusion of Si on Si Particles in the Matrix ,000 10, ,000 t c (s) Figure 9. Critical times for the three quench precipitates. Hold times and temperatures for specimens analyzed metallographically are also indicated. 22

23 Predicted Y (MPa) M. Tiryakioǧlu, R.T. Shuey: Metall. Mater. Trans. B, v. 38, pp , CWQ 450 C 400 C 350 C 300 C DQ R 2 = Experimental Y (MPa) Figure 10. Experimental versus predicted yield strength for D357 specimens (CWQ: Cold water quench, DQ: delayed quench) 23

24 o C 450 C 400 C 350 C 300 C Y (MPa) o C o C 400 o C ,000 10, ,000 1,000,000 Hold Time (s) Figure 11. The predicted effect of hold time at various temperatures on yield strength of D357 specimens with experimental data indicated. 24

25 T ( o C) M. Tiryakioǧlu, R.T. Shuey: Metall. Mater. Trans. B, v. 38, pp , % % ,000 10,000 time (s) Figure % and 90% iso-yield strength contours. 25

26 T ( o C) M. Tiryakioǧlu, R.T. Shuey: Metall. Mater. Trans. B, v. 38, pp , (0.65Si-0.89Mg-0.23Cu) (0.92Si-0.59Mg) D357 (7.11 Si-0.61Mg) ,000 10,000 time (s) Figure % iso-yield strength contours in three Al-Mg 2 Si alloys. 26

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