CONTROL OF PHASE TRANSFORMATION DURING HEAT TREATMENTS BASED ON DSC EXPERIMENTS

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1 Proceedings of the International Conference on Solid-Solid Phase Transformations in Inorganic Materials 2015 Edited by: Matthias Militzer, Gianluigi Botton, Long-Qing Chen, James Howe, Chadwick Sinclair, and Hatem Zurob CONTROL OF PHASE TRANSFORMATION DURING HEAT TREATMENTS BASED ON DSC EXPERIMENTS Philipp Schumacher 1,2, Stefan Pogatscher 3, Marco J. Starink 4, Olaf Keßler 1, Christoph Schick 2, Volker Mohles 5 and Benjamin Milkereit 1,2,4 1 University of Rostock; Chair of Materials Science; Germany 2 University of Rostock; Polymer Physics Group; Germany 3 ETH Zurich; Laboratory of Metal Physics and Technology; Switzerland 4 University of Southampton; Engineering and the Environment; United Kingdom 5 RWTH Aachen University; Institute of Physical Metallurgy and Metal Physics; Germany Keywords: Al-Si alloys, DSC analysis, Quench-induced precipitates, Microstructure control, Mechanical properties testing Abstract Phase transformation in Al-Si alloys during cooling from solution annealing was investigated with advanced DSC techniques in a wide cooling rate range (2 K/s K/s). Microstructural analyses show that Si particles of different shape and size can precipitate during cooling. The dependence of precipitation enthalpy on cooling rate and temperature is modelled, providing a consistent physical description of precipitate volume fraction and solute Si amount during cooling. This allows control of precipitation via precise heat treatments and thus allows generation of welldefined microstructural states. Thereby, samples having an equal amount of solute Si but different precipitation states could be tested in order to investigate their influence on the mechanical behaviour. A big advantage of this method is that the strengthening contribution of precipitates can be determined without the need to assume any - potentially inaccurate - superposition law between particle and solute strengthening. Introduction In general, strength of commercial aluminium alloys originates from a combined effect of several factors such as dislocation density, solid solution state, grain size and particles. As the mechanical properties are controlled by all aspects of microstructure, it is difficult to separate strengthening contributions of single factors [1]. Furthermore, since microstructure evolution is strongly related to processing conditions of the material, comparisons of properties between different material conditions are often not suitable to draw conclusions on active strengthening mechanisms. However, reliable experimental data of precise microstructural states is needed for validation of plasticity models for through-process modelling, which are based on microstructural approaches and hence distinguish different dislocation interaction mechanisms. It is therefore necessary to tailor well-defined microstructural states that can be mechanically tested. An appropriate method for this is the application of specific heat treatments, as they allow adjustment of the microstructure of aluminium alloys to a great extent. The cooling procedure subsequent to solution annealing significantly impacts the microstructure of the alloys and can be utilised to modify solute levels as well as types and volume fractions of precipitated particles. Hence, by adjustment of cooling conditions microstructural states can be tailored that allow investigation of strengthening contributions of solute atoms and secondary phases. The main difficulty in doing this is the determination of appropriate heat treatment parameters, because well- 699

2 defined material conditions can only be achieved if the occurring phase transformations during heat treating are known. It was shown in several publications (e.g. [2 5]) that differential scanning calorimetry (DSC) is a powerful technique to study phase transformations in aluminium alloys. The measured precipitation heat (enthalpy change) upon cooling is directly proportional to the volume fraction of precipitated particles [3,6]. Hence, precise determination of formation enthalpies for the entire cooling rate range of physical interest provides input for a physically-based model that allows a calculation of remaining solute amount in dependence of cooling rate and temperature. Information from DSC measurements in a wide dynamic range are therefore of great importance to set heat treatment parameters for the control of microstructure, i.e. the microchemical composition of the matrix and the corresponding precipitation state of aluminium alloys. In the present work a recently introduced experimental approach [7] on how to study the correlation between microstructural states of non-equilibrium alloy conditions and their mechanical behaviour is applied. Materials and ethods Experimental work was carried out on binary Al-Si alloys of high purity, which allow a detailed study on strengthening contributions of Si solutes and quench-induced Si precipitates on mechanical properties. Chemical compositions of the alloys are shown in Table I. Both alloys were processed equally prior to all experiments. The initial conditions of the materials and sampling procedures for all experimental techniques are described in Ref. [7]. Table I. Chemical composition of investigated alloys obtained by optical emission spectroscopy (OES) analysis. mass fractions Si Mg Fe Cu Mn Cr Zn Ti (%) (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) Al-0.26Si 0.26 <1 <5 4 <3 <3-20 Al-0.72Si 0.72 <1 <5 <1 <3 < Precipitation in both Al-Si alloys during cooling from solution annealing (540 C, 20 min) have been studied with DSC in a wide cooling rate range. The principle measurement and evaluation methods employed for these experiments are described in Ref. [2]. Recently, direct cooling experiments with rates down to 10-3 K/s have been successfully implemented using a heat flux DSC Setaram C600 equipped with a high precision 3D Calvet sensor. Furthermore, a new indirect DSC reheating method has been introduced [7], which allows investigations on enthalpy changes caused by precipitation reactions at cooling rates down to 10-4 K/s (possessing potential of even slower cooling). Figure 1 gives an overview on the applied methods and the utilised devices. In addition, cooling durations from solution annealing to 20 C are indicated for several rates. Figure 1. Applied methods and devices for DSC investigations on precipitation upon cooling in a wide dynamic range. 700

3 For the characterisation of quench-induced precipitates, metallographic investigations with light optical microscopy (LOM) were carried out on selected cooling states. From DSC results a consistent physical description of remaining amount of solute Si in dependence of cooling rate and temperature was derived. As only the diamond cubic Si-rich equilibrium phase ( = nm) is assumed to be precipitated during cooling from solid solution [8], the model additionally allows to calculate volume fractions of quench-induced Si precipitates after distinct heat treatments. The model was validated in Ref. [7] using atom probe tomography (APT) and computational image analysis. Flow curves of various microstructural conditions were determined by compression tests on cylindrical samples (Ø5 mm x 10 mm). All mechanical tests were performed in a quenching and deformation dilatometer Bähr DIL 805 A/D at a temperature of 30 C, while strain rates of 10-1 s -1 and 10-3 s -1 were applied. Well-defined heat treatments were executed before testing. A necessary precondition for the intended comparisons was that all strengthening factors not under investigation had to contribute equally in the examined material states. Grain size effects were negligible in the investigated alloys, because both alloys showed very coarse grains with similar sizes of about 450 m after heat treatments. Due to the high purity of the alloys, influences from any constituent particles and dispersoids could also be neglected. Hence, for the investigation of strength contribution of Si solutes, solely an equal precipitation state in both alloys was required. On the contrary, an identical solid solution state was essential to investigate strengthening effects of precipitates formed during cooling. For both alloys precipitate-free conditions were generated to investigate the influence of solute Si. Furthermore, equally processed high purity Al was tested for reference purposes. Tailoring of these states was executed directly in the dilatometer. Annealing was performed in a vacuum atmosphere. Apart from a precipitate-free condition, defined states of Al-0.72Si with a remaining solute amount of 0.26 ma.% Si and quench-induced Si precipitates of different morphology were generated. Mechanical properties of these states were compared with Al-0.26Si in complete solid solution. Thus, samples having an equal amount of solute Si but different precipitation states were analysed. Appropriate heat treatment parameters to tailor these microstructural states were derived from the introduced model. Heat treatments were carried out in a chamber furnace with cascade control and samples were sealed in quartz glass under vacuum prior to heat treatment in order to prevent surface reactions and thus changes in chemical composition. Before the samples were mechanically tested in the dilatometer, they were stored at -82 C in order to supress further potential precipitation reactions. Precipitation upon cooling Results and iscussion Excess specific heat capacity curves of Al-0.72Si during cooling with two different rates after solution annealing (540 C, 20 min) are shown in Figure 2. Exothermic precipitation reactions are shown by deviations exceeding the zero level, which is given by a dashed line for both curves. Curves over a much larger cooling rate range are presented in Ref. [7]. However, data displayed in Figure 2 suffice to expresses the fundamental precipitation behaviour of the alloy. Intensity of precipitation decreases with increasing cooling rate. This can be explained by the suppression of diffusion processes. Interestingly, two reaction peaks for alloy Al-0.72Si within the analysed cooling rate range have been observed. This may seem surprising, as only diamond cubic Si-rich phase is likely to occur [8,9]. For a cooling rate of K/s, precipitation reactions in Al-0.72Si start at about 475 C, which is very close to the solvus temperature of this binary alloy system [10]. The high temperature precipitation peak, which is dominating at this cooling condition, is 701

4 Figure 2. DSC cooling curves of alloy Al-0.72Si after solution annealing. Marked temperatures will be explained below. suppressed significantly in its intensity and shifted towards lower temperatures as the cooling rate increases. Concurrently, another peak at lower temperatures rises to a maximum. Such precipitation behaviour is also shown by commercial alloys of different alloying systems (e.g. [3,5]). Further increase in cooling rate results in a decrease of this low temperature peak area as well. The upper critical cooling rate at which the complete solid solution is retained during cooling was found to be about 1 K/s for Al-0.72Si. As the upper critical cooling rate is much lower for Al-0.26Si, this rate also results in a complete suppression of precipitation reactions in this alloy. Different DSC reaction peaks during cooling correspond to the formation of different types of precipitates [3,5]. However, no precipitation sequence is known for binary Al-Si alloys. Since at least two exothermic reactions were detected for alloy Al-0.72Si, metallographic studies were carried out to analyse the development of precipitate formation by a step quenching procedure introduced in Ref. [3]. Samples were cooled with K/s and 0.01 K/s to specific temperatures, before they were overcritically cooled to maintain the precipitation state of interest. Chosen temperatures are marked in Figure 2. The microstructural development of Al-0.72Si during cooling with the investigated rates is shown in Figure 3. LOM images revealed that precipitation of particles with different morphologies seem to correspond to the reaction peaks measured by DSC. For both rates particles are mainly precipitated inside aluminium solid solution grains. When the alloy is cooled with K/s, no precipitates are visible at 490 C. This coincides with the corresponding DSC curve. No reactions were detected above this temperature. Coarse polygonal particles (dimensions some m) with low Figure 3. Development of precipitate formation in Al-0.72Si upon cooling with a) K/s and b) 0.01 K/s after solution annealing to specific temperatures followed by overcritical cooling. 702

5 aspect ratios have been formed at 440 C, where the intensity maximum of the high temperature reaction peak is reached (Figure 2). With a further decrease of temperature the precipitates grow and the volume fraction increases. For a cooling rate of 0.01 K/s the maximum amount of heat is released by low temperature reactions. However, high temperature reactions are not fully suppressed at that rate. At a temperature of 435 C only polygonal-shaped precipitates with much smaller dimensions compared to the slower cooling condition can be perceived, which correspond to high temperature reactions taking place in this temperature region (Figure 2). When a temperature of 395 C is reached (low temperature reaction peak maximum) these particles have grown and first rods or platelets appear. With further decrease of temperature the polygonal-shaped particles seem to retain their form and size. This indicates that these particles form only at high temperatures (high temperature precipitates). On the contrary, amount and size of rod/plateletshaped particles increases after passing the intensity maximum of the low temperature reactions (low temperature precipitates). The rod/platelet-shaped particles have high aspect ratios with lengths up to 10 m and widths of some 100 nm. Besides the knowledge on precipitated particle types, a model to calculate the remaining amount of solute Si in dependence of cooling rate and temperature is needed. It was shown in Ref. [7] that solute Si content in Al-Si alloys can be estimated from the following equation, in which parameter csi total corresponds to the overall content of Si in the material: (1) The enthalpy change HT c ( ) caused by precipitation reactions up to temperature Tc, to which the alloy has already been cooled with cooling rate, can be computed by integration of DSC excess specific heat capacity (excess cp) curves. The solution annealing temperature Tan is taken as upper bound of integration: (2) To calculate the amount of solute Si with Eq. (1), the saturation value of enthalpy change HRT sat after cooling to room temperature must be known. This value is reached if complete precipitation of Si takes place during cooling. As solubility of Si in the Al-matrix is negligible, HRT sat is a) b) Figure 4. Change in enthalpy upon precipitation in Al-0.72Si: a) Enthalpy change HRT as a function of cooling rate obtained with direct DSC cooling experiments or indirect DSC reheating experiments and model fit [7]. b) Enthalpy change HT c for two cooling rates as a function of temperature. 703

6 proportional to the Si content of the alloy. An appropriate model based on diffusion controlled reactions [11,12] to compute enthalpy change after cooling with different rates to room temperature was introduced in Ref. [7]. A sufficient amount of data is needed to achieve a model of high quality, which underlines the significance of reliable DSC experiments in a wide cooling rate range. Especially data for very low cooling rates is needed to estimate HRT sat with certainty. Figure 4 a shows enthalpy change values after cooling to room temperature HRT obtained from direct and indirect DSC measurements as a function of cooling rate. Error bars express uncertainties of 10% resulting from peak-area determination of DSC data. Furthermore, the resultant model curve is displayed. The adapted model fits the data of Al-0.72Si very well for cooling rates covering more than five orders of magnitude. The necessary saturation value HRT sat for calculation of solute Si content was estimated to be 11.3 J/g. Tailoring well-defined microstructural states The introduced methods allow the adjustment of precise material conditions. Table II gives an overview on adjusted states for mechanical testing and the corresponding heat treatment parameters (cooling conditions after solution annealing). Table II. Material conditions and applied heat treatment parameters for investigations on strength contribution of Si solutes and quench-induced Si precipitates. alloy solute mass cooling condition after predominant particle type fraction Si solution annealing Al-0.72Si 0.72 precipitate-free = 1 K/s, T c =30 C Al-0.26Si precipitate-free = 1 K/s, T c = 30 C Al-0.72Si Al-0.72Si 0.26 high temperature precipitates - volume fraction: 0.5 % - shape: polygonal - size: 3-5 m low temperature precipitates - volume fraction: 0.5 % - shape: rods/platelets - length/width: <10 m/<1 m = K/s, T c =400 C subsequent water quenching = 0.01 K/s, T c =30 C Al precipitate-free = 1 K/s, T c = 30 C Heat treatment parameters were determined from DSC data. Precipitate-free conditions of both alloys were adjusted by overcritical cooling with 1 K/s from solution annealing. Parameters to tailor the other states were derived as follows. At first, cooling rates at which predominantly different particle types are precipitated were chosen. Appropriate conditions can be adjusted with the presented rates of K/s and 0.01 K/s, as they cause very different precipitation reactions resulting in dissimilar microstructures. With application of Eq. (1) and Eq. (2) specific temperatures Tc, at which a Si mass fraction of 0.26 % remains in solution, were then determined for both cooling rates. The procedure for the determination of heat treatment parameter Tc is visualised in Figure 4 b. Plotted are enthalpy changes HT c as a function of temperature calculated with Eq. (2) for both cooling rates. With decreasing temperature, HT c increases due to precipitation reactions. When reactions are completed, a constant value is attained, which is dependent on the cooling rate. With knowledge of saturation value HRT sat the necessary enthalpy change for a defined solute content can be calculated with Eq. (1). Temperature Tc to which samples need to be cooled can then directly 704

7 be derived from the integrated DSC curves. For a better understanding the solute mass fraction Si for all possible values of enthalpy change is given as second ordinate in the diagram. For a Si mass fraction of 0.26 % to remain in solution, an enthalpy change of HT c 7.2 J/g is necessary. If precipitation reactions are still detectable (enthalpy change constant) at temperature Tc samples have to be cooled overcritically as soon as Tc is reached to maintain the precipitation state. Mechanical testing of well-defined microstructural states The precipitate-free conditions of both alloys show different flow characteristics (Figure 5 a). A higher amount of solute Si enhances the yield strength of Al-Si alloys. However, comparison to pure Al shows that the strengthening effect of Si is rather small. While a higher solute Si amount results only in a slight increase in the resistance to plastic deformation, work hardening is strongly influenced by the Si content. Increased work hardening compared to Al was observed in both Al-Si alloys. Hardening was almost linear up to a total strain of about 0.04 for the two conditions. Further increase of strain leads to non-linear stress-strain curves, whereby Al-0.72Si shows much higher work hardening potential. By comparison of conditions having an equal solute Si content of 0.26 ma.% but different precipitation states, identical yield strengths were observed for all conditions (Figure 5 b). This indicates that both precipitate types formed during cooling have no impact on initial resistance to plastic deformation. However, influence of precipitated particles becomes significant as the strain increases. Higher strain hardening at the initial stage of plastic deformation was observed for samples containing precipitates. Precipitate-free Al-0.26Si thus showed a more linear hardening, while a parabolic hardening was observed for both states containing quench-induced precipitates. Interestingly, work hardening is much stronger in Al-0.72Si samples with small rod/platelet-like particles, which are formed at lower temperatures. Strength of this condition even exceeds the precipitate-free Al-0.72Si state (Figure 5 a) in the strain range between 0.01 and The results clearly show that quench-induced precipitates in Al-Si can influence the mechanical properties despite their relatively large sizes. While coarse high temperature precipitates with low aspect ratios seem to have a very low strength contribution, the smaller rod/platelet-shaped precipitates formed at lower temperatures clearly influence work hardening. Mechanical tests with a strain rate of 10-3 s -1 lead to same results, indicating that flow characteristics of the investigated conditions are not strain rate dependent at low temperatures. Figure 5. Comparison of stress-strain curves at 30 C to investigate strengthening contribution of a) solute Si and b) quench-induced precipitates in Al-Si alloys (strain rate: 10-1 s -1 ). 705

8 Summary An approach to explore strengthening contribution of solutes and quenched-induced precipitates has been introduced on binary Al-Si alloys. Well-defined material conditions were adjusted by heat treatments and mechanically tested. Strengthening contributions could be discussed without the need to assume any superposition law between particle and solute strengthening. Appropriate heat treatment parameters were derived from extensive DSC analyses on occurring precipitation reactions. It was shown that form and size of quench-induced Si particles depend on the cooling rate and temperature range in which they are precipitated. The mechanical tests showed that higher amount of solute Si slightly increases yield strength and leads to much higher work hardening potential. On the other hand, quench-induced Si precipitates are too coarse to contribute to the alloy s yield strength. The various precipitate types still appear to have different influences on hardening. While coarse polygonal particles with low aspect ratios formed at high temperatures seem to have a weak strengthening contribution, smaller rod/plateletshaped precipitates cause much stronger hardening with increasing amount of strain. Acknowledgements The authors gratefully acknowledge funding of this work by Deutsche Forschungsgemeinschaft (MO974/4-1 and MI1731/1-1). Besides, part of this work was supported by a fellowship within the Postdoc-Program of the German Academic Exchange Service (DAAD) for B. Milkereit. References [1] Ø. Ryen et al., "Strengthening mechanisms in solid solution aluminum alloys", Metall Mat Trans A Phys Metall Mat Sci, 37 (2006), [2] B. Milkereit, O. Kessler, C. Schick, "Recording of continuous cooling precipitation diagrams of aluminium alloys", Thermochim Acta, 492 (2009), [3] B. Milkereit et al., "Continuous cooling precipitation diagrams of Al-Mg-Si alloys", Mater Sci Eng A, 550 (2012), [4] D. Zohrabyan et al., "Continuous cooling precipitation diagram of high alloyed Al-Zn-Mg- Cu 7049A alloy", Trans Nonferrous Met Soc China, 24 (2014), [5] Y. Zhang et al., "Development of continuous cooling precipitation diagrams for aluminium alloys AA7150 and AA7020", J Alloys Compd, 584 (2014), [6] M.J. Starink, "Analysis of aluminium based alloys by calorimetry: Quantitative analysis of reactions and reaction kinetics", Int Mater Rev, 49 (2004), [7] P. Schumacher et al., "Quench-induced precipitates in Al Si alloys: Calorimetric determination of solute content and characterisation of microstructure", Thermochim Acta, 602 (2015), [8] I.J. Polmear, Light alloys (Oxford: Elsevier/Butterworth-Heinemann, 2006). [9] M.J. Starink, A.-M. Zahra, "Kinetics of isothermal and non-isothermal precipitation in an Al - 6 at.% Si alloy", Philos Mag A, 77 (1998), [10] J.L. Murray, A.J. McAlister, "The Al-Si (Aluminum-Silicon) system", Bull Alloy Phase Diagr, 5 (1984), [11] M.J. Starink, A.-M. Zahra, "An analysis method for nucleation and growth controlled reactions at constant heating rate", Thermochim Acta, 292 (1997), [12] M.J. Starink, "On the meaning of the impingement parameter in kinetic equations for nucleation and growth reactions", J Mater Sci, 36 (2001),

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