Impact of Mn on the precipitate structure and creep resistance of Ca containing magnesium alloys

Materials Science & Engineering A 761 (2019) 137964

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Impact of Mn on the precipitate structure and creep resistance of Ca containing magnesium alloys

T

Steffen Lamm*, Dorothea Matschkal, Mathias Göken, Peter Felfer Institute for General Materials Properties, Department Materials Sciences, Friedrich-Alexander University Erlangen-Nuremberg, Germany

A R T I C LE I N FO

A B S T R A C T

Keywords: Magnesium alloys Creep Clusters Precipitate structure

The reason for the enhanced creep resistance of Ca-modified magnesium alloys is investigated based on their microstructure and precipitate morphology. It is known from literature that the creep resistance of commercial AZ91 is improved by the addition of Calcium through the formation of a skeleton-like intermetallic phase (IP) at the grain boundaries. Compared to AZ91 + Ca, the alloy MRI230D, which is a Mg–Al–Ca–Sn based alloy with similar Ca content, shows a significantly better creep resistance. This is probably caused by small precipitates in the α-Mg grain interior that were found in earlier TEM investigations of as-cast MRI230D alloy. For this work, we have investigated the evolution of the precipitate structure from as-cast to crept conditions via atom probe tomography (APT). Based on the APT data, an intensive cluster analysis for all deformation states was performed using Voronoi cluster search algorithm. Cluster analysis revealed that Ca atoms show an especially high grade of clustering in as-cast conditions. This changes in the crept states, where co-clustering of Ca and Mn atoms was observed. Further analysis revealed that the number density of clusters is generally higher in MRI230D than in AXZ931 and clusters tend to be smaller and enriched with Mn. All of this led to the conclusion that a minimum amount of Mn is necessary for a stabilization of the clusters in the Mg matrix, leading to an improvement of the creep resistance.

1. Introduction

relatively harmless intermetallic compounds [3,4]. The reason for their relatively poor high temperature strength and creep resistance has been attributed to the presence of the β-Mg17Al12 intermetallic phase, which easily coarsens at elevated temperatures [5,6]. It is also reported that the occurrence of dynamic discontinuous grain boundary precipitates of Mg17Al12 facilitates grain boundary sliding, due to its migration into the neighboring grains during growth. As a result, the general strategy adopted to improving the creep resistance of Mg alloys is to have thermally stable intermetallic phases at grain boundaries [5–8]. It is reported in literature [2,9–11] that the creep behavior of Mg–Al alloys can be improved by the addition of elements such as Sb, Ca, Sr, Si and RE (rare earth elements), which lead to the formation of thermally stable and morphologically favorable intermetallic phases. The good creep properties in AE42, for example, derive mostly from the suppression of the Mg17Al12-phase and instead formation of RE containing precipitates [10]. RE elements can strongly increase the creep resistance at elevated temperatures, but those elements are also comparably expensive. A more cost effective element to increase the creep strength is Ca. Ca can increase the creep strength of RE-magnesium alloys like Mg–4Al-RE [12], but also in commercial non-RE Mg-alloys like AM60 [11] or AZ61 [13]. Strengthening of these alloys by calcium

The low density and high specific strength of magnesium alloys make them very suitable for lightweight applications in the automotive industry, where Mg alloys are widely used in instrument panel beams or steering components. Mg alloys have also become more interesting for use in powertrain components, in order to reduce vehicle curb weight. In these applications, the magnesium parts have to withstand temperatures in the range of 150 °C–300 °C. Conventional alloys, like AZ91, show an unfortunately poor strength at these elevated temperatures, which led to the development of more creep resistant variants like AXZ931 (Mg–Al–Zn–Ca–Mn) and MRI230D (Mg–Al–Zn–Ca–Mn–Sr–Sn), while still retaining suitability for die casting. The most common Mg alloys used for die casting are based on the Mg–Al–Zn or Mg–Al–Mnsystem. Alloys with or without the addition of zinc, like AZ91 (Mg–9Al–1Zn) [1] or AM60 (Mg–6Al-0.3Mn) [2], are widely used because of their good mechanical properties at room temperature, good castability and low cost. They also show good corrosion resistance, which is due to the presence of Mn and Zn. Zn additionally improves the room temperature strength. Mn increases the corrosion resistance in these alloys by binding Fe and other heavy-metal elements into

*

Corresponding author. E-mail address: steff[email protected] (S. Lamm).

https://doi.org/10.1016/j.msea.2019.05.094 Received 11 April 2019; Received in revised form 20 May 2019; Accepted 25 May 2019 Available online 13 June 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.

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AZ91 [20]. In this work, we will be focusing on understanding the creep behavior of the two alloys AXZ931 and MRI230D and the different mechanisms of precipitation evolution during creep and ageing at application relevant temperatures (200 °C). We use atom probe tomography (APT) to determine the precipitate chemistry and its evolution during creep deformation and after ageing at 200 °C without mechanical stresses. In Fig. 2 creep curves of those alloys are shown. These creep tests were performed at 200 °C with loads of 80 and 100 MPa for the AXZ931 and 100 MPa for the MRI230D alloy. In Fig. 2a, all creep curves show a quick reduction of the strain rate at the beginning, slowing down to a creep minimum, followed by an increase of the strain rate. If we compare the tests with the same load of 100 MPa, it is clearly visible that the minimum creep rate (MCR) of AXZ931 is significantly higher than the minimum creep rate of MRI230D, i.e. AXZ931 is softening much faster than MRI230D. This is a further evidence that the creep rate is governed by the stability of the precipitates, as the IP network remains intact during small deformations. At the same load, the testing time and in conjunction, time for coarsening of the precipitates in AXZ931 is much shorter than that for MRI230D (Fig. 2b), which could have an influence on the creep response. To study this in more detail, we investigated the microstructure of MRI230D deformed at 100 MPa with samples of AXZ931 deformed at 80 MPa, because these samples have comparable long creep times (Fig. 2b).

is achieved by the formation of the Al2Ca phase and an accompanied reduction of the Mg17Al12 phase contents, as presented in ref. [13]. In Ref. [14], for example, an extruded Mg-3.5Al-3.3Ca-0.4Mn alloy exhibited improved tensile strength because of nanoscale plate-like Al–Ca precipitates and spherical Al–Ca–Mn precipitates, along with strong basal texture. Another strengthening effect due to the addition of calcium originates from solid solution strengthening of the Mg matrix by Ca and the particle strengthening effects of the thermally stable Al2Ca phase which is formed as an interconnected network [11]. A similar effect was observed in Ref. [15] where an addition up to 5 wt.-% calcium to AZ91 lead to an improvement of the minimum creep rate. This improvement in creep properties of these Ca-added AZ91 alloys is also due to the formation of a skeleton like intermetallic phase (IP) at the grain boundaries. Beside the attempt to improve the creep resistance of established alloys, a creep resistant alloy named MRI230D was developed to operate at temperatures up to 190 °C [16]. MRI230D includes additional elements like Sr and Sn. According to Ref. [4], Sr addition up to 0.5 wt.-% leads to a modification of the intermetallic phases and further improvement of the creep resistance. Sn is added to eliminate sticking to the die. Typical loads at elevated temperatures are between 50 MPa and 100 MPa, according to Refs. [10,17,18], if applications in engine near parts like gearbox housings and crankcases are considered. In the following investigations, 100 MPa is the lower limit for the chosen creep stress for MRI230D. This is to limit the time necessary for the creep tests. In earlier work [15], the creep properties of MRI230D (~Mg-6.8Al-1.9Ca-0.4Mn-0.3Zn-0.5Sr-1.1Sn in wt.-%) were compared to a series of alloys based on AZ91 with additions of calcium up to 5 wt.%. The limit of the Ca addition was set to 5 wt.-% because further increasing the Ca content lead to a decreased elongation to failure [19]. It was stated before that Ca addition leads to the formation of a skeleton like intermetallic phase (IP) at the grain boundaries. Investigations of this AZ91 alloys series, with additions of up to 5 wt.-%, showed that the interconnectivity of the IP increases with increasing Ca addition. As a result, the IP network was identified to be the dominating contributor to the improved creep resistance. A micrograph of the alloy with 3 wt.-% Ca (AXZ931) is shown in Fig. 1a. Interestingly, the MRI230D alloy (Fig. 1b) exhibits superior creep properties compared to AXZ931, although the interconnected IP network is quite similar. Also compared to the alloy with 5 wt.-% Ca addition (AXZ951), where the IP interconnectivity is higher than in MRI230D, still MRI230D shows a lower minimum creep rate (MCR) with 3*10−8 s−1 compared to ~7*10−8 s−1 of AXZ951 [15]. Further, TEM investigations revealed nanometer size, precipitate like structures in the MRI230D grain interior, while there were no such features identified in the grain interior of the Ca alloyed AZ91 series. This is a strong indication that small scale precipitates in the grain interior are largely responsible for the improvements in creep strength. These assumptions were confirmed by nanoindentation of the grain interior, which revealed an increased matrix hardness for MRI230D compared to

2. Materials and methods 2.1. Materials In the present work, the alloys AXZ931 (AZ91 with 3 wt.-% Ca addition) and MRI230D that were already characterized in Ref. [15] were chosen to analyze the precipitate structure via APT. The production of the alloys is described in detail in Ref. [15]. The alloys were thixomolded by Neue Materialien Fürth GmbH on a Japan Steel Works 220 t machine at a casting temperature of 878 K. Casted parts with 6 mm wall thickness were chosen. Table 1 summarizes the chemical composition of the investigated alloys chosen for the present work as determined by glow discharge optical emission spectroscopy (GDOES) using a GD-PROFILER (HORIBA Jobin Yvon).To analyze the precipitate evolution, APT analysis was performed on samples in as-cast condition, after creep at 200 °C and in aged conditions at 200 °C. Testing time for the creep samples was 232 h at 80 MPa for AXZ931 and 257 h at 100 MPa for MRI230D (see also Fig. 2 b). For comparison, samples were aged for 250 h at 200 °C to compare thermal and thermomechanical influences on the resulting microstructures. Additionally, a sample of MRI230D was aged for 500 h in order to look at long term evolution of the more stable precipitate microstructure of MRI230D.

Fig. 1. SEM micrographs of the (a) thixomolded AXZ931 and (b) MRI230D alloys. Images show the α-Mg matrix with IP at the grain boundaries. 2

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Fig. 2. The strain rate-strain curves (a) of compression creep tests at 200 °C and a load of 80 and 100 MPa, show the much lower minimum creep rate of MRI230D compared to AXZ931. The strain-time curves in (b) clearly show the much shorter creep times for AXZ931, at the same load of 100 MPa.

treatment) are shown in Fig. 3, with data from the MRI230D alloy on the left and AXZ931 on the right side. All reconstructions show volumes of the grain interior regions. In every condition, 1 to 3 atom probe measurements with more than 25 million atoms were investigated. In the reconstructions in Fig. 3 aluminum atoms are displayed in light blue to outline the captured volumes. For visual clarity, only ~1% of all detected Al atoms are displayed. The regions with enriched atoms of one species are visualized by the 1%-isoconcentration surfaces of Ca (blue) and Mn (red), generated using the marching cubes algorithm [24] on a voxelisation with 1 × 1x1 nm3 voxel size. These isosurfaces can act as a rough outline for the fine precipitates. The fact that many of them appear to have a common alignment that is not related to the APT analysis direction implies the presence of an orientation relationship with the parent Mg crystal. Since only one variant of this relationship is observed, it can be assumed that larger facets of the precipitates are lying on the (0001)planes of the Mg crystal, like it was observed in Ref. [25], for an Mg–5Al–3Ca-0.15Sr (wt.-%) alloy where Al2Ca precipitates were found on the basal planes of the α–Mg matrix phase. The Al content measured with APT is around 2 to 3 atomic percent, which is in good accordance with the maximum solubility of ~2.8 at.-% Al in Mg at 200 °C [26]. However, in Ref. [19] the authors only calculated 2.25 wt.-% Al (~2 at.% when converted) in solid solution of AXZ931 in electron probe

2.2. Methods APT experiments were carried out in a Cameca LEAP 4000X HR with a reflectron detector system offering ~37% detection efficiency. The samples were prepared via FIB lift-out in a FEI Helios 660 and a Zeiss Crossbeam 540. While the lift-out process was done with a 30 kV Ga-beam, the shaping of the samples as soon as they were smaller than 1 μm was carried out under reduced acceleration voltage of 10 kV. This eliminates spurious ion irradiation damage within the investigated volumes caused by ion channeling and with it any unwanted alterations to the precipitates. For the final milling step, the ion energy was reduced to 5 kV. A more precise description of the lift-out procedures are given in Refs. [22,23]. All APT measurements were performed in voltage pulsed mode at a temperature range of 40–45 K. The pulse-fraction (ratio of pulse voltage to bias voltage) was kept between 15 and 20 % with the bias voltage adjusted to an average detection rate of 1 atom per 100 pulses. 3. Results A first approach for visualization of APT data is to only plot the nonhomogeneously distributed atoms in a reconstruction. The APT reconstructions of all investigated alloy conditions (as-cast, crept, thermal

Fig. 3. APT reconstructions of all investigated alloy conditions. Reconstructions show the 1 at.-%-isosurfaces for calcium (blue) and manganese (red). Aluminum atoms (light blue) are displayed to point out the size and shape of the measured APT samples. The scalebar applies for all reconstructions. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

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Table 1 Chemical Composition of the investigated alloys in wt.-%, measured via glow discharge optical emission spectroscopy (GDOES) [15,21] and converted values in at.-%.

wt.-% at.-% wt.-% at.-%

Alloy

Mg

Al

Zn

Ca

Mn

Sr

Sn

MRI230D MRI230D AXZ931 AXZ931

89.16 91,86 88.04 90.45

6.76 6,28 8.02 7.4

0.27 0,1 0.55 0.21

1.94 1,21 2.96 1.84

0.38 0,17 0.22 0.1

0.48 0,14 – –

1.1 0,23 – –

microanalysis (EPMA) measurements. The Zn content measured with APT, however, is in good accordance with [19], being ~0.1 at.-% (~0.28 wt%). The measured Ca content, with 0.03–0.07 at-%, is strongly below the nominal composition of the alloys (see Table 1), which is an indicator for high segregation of Ca atoms to the IPs. The general trend that Ca is found in the eutectic IPs is also observed in Ref. [19]. The Mn content, as measured by APT in our AXZ931 alloy, is in the range of 0.04–0.05 at.-% and in the range of 0.04–0.08 at.-% in the MRI230D measurements. This means that the Mn atoms also segregate to the IPs, although the segregation is not as strong as for Calcium. The amount of Zn measured in the α-Mg grain interior of AXZ931 is only a factor of two lower than the nominal composition (Table 1), Sn and Sr concentrations are below 0.07 at.-% and are left out in the visualization of Fig. 3. Fig. 3a and b show APT reconstructions of the as cast-conditions, including 1%-Ca isoconcentration surfaces (blue), showing the aforementioned orientation relationship with the Mg crystal. It should be mentioned that no enrichment of Mn was observed in as-cast conditions. Fig. 3c and d depict reconstructions of samples after creep deformation at 200 °C, together with 1% Ca and 1% Mn isosurfaces. Here, a distinct change in precipitation structure over the as-cast condition is evidenced by the co-location of Mn with Ca for both MRI230D and AXZ931. From the number density of visually discernible precipitates in Fig. 3c, it can already qualitatively be stated that the number density and volume fraction of precipitates has significantly increased over the as-cast condition. This is not as distinct in the case for AXZ931 in Fig. 3d. In Fig. 3e and f the alloy conditions after aging for 250 h at 200 °C are shown in the same fashion as for the crept microstructures in Fig. 3c and d. This data corresponds to an aging time similar to the creep deformation applied to the samples in the above figures. While qualitatively the change in structure is at least limited in the case of MRI230D, the size and distribution of the Ca-isosurfaces in AXZ931 (Fig. 3f) is significantly coarser. For the MRI230D alloy, we prepared a further aged condition (500 h, 200 °C) to gauge the thermal stability of the precipitate structure. This is shown in Fig. 3g. This structure is similar to the crept state and the 250 h aged conditions. Only the number density and size of the Ca-rich precipitates appears to have grown. This will be discussed later in detail. Before that a first approximation of the overall volume-fraction of the precipitates can be obtained from the meshes of the Ca + Mn isosurfaces, which were generated by the marching cubes algorithm [25]. These results of the volume-fraction calculations are shown in Fig. 4. It is observed that the volume fractions of the isosurfaces are below 1% in the as-cast states of both alloys. After creep deformation, the volume fraction determined by the isosurfaces more than doubles to over 2 vol.-% for both alloys. However the aged conditions of the alloys show different trends in obtained isosurface volume fractions. While the volume fraction slightly increases to ~2.4 vol.-% in the case of the aged AXZ931, the overall amount of precipitation seems to be reduced for the aged MRI230D (~1.7 vol.-% precipitation). The 500 h aged sample of MRI230D shows a precipitation volume fraction of ~1.9 vol.-%, which is slightly higher than after 250 h ageing, but still lower than after creep deformation. This indicates that the full precipitation potential of the alloy has not been reached even after 500 h of aging at this

Fig. 4. Analysis of the volume fraction of the combined Ca + Mn isosurfaces, displayed in Fig. 3.

temperature. Based on the above analyses, we assessed the chemical makeup of the precipitates using proximity histograms [27]. In proximity histograms, the concentration along a distance to a reference object is plotted. Since the investigated precipitates are very small, we chose the distance to the precipitate center as the relevant coordinate. This minimizes interference from local magnification effects [28] around precipitates in the atom probe. The precipitate centers were obtained by a preliminary search for clusters of Al, Ca and Mn atoms. This search was performed using a Voronoi cluster search algorithm [29] on clusters containing Al, Ca and Mn atoms, placing them at the center of mass of the individual detected clusters. It should be noted that for this reason, the proximity histograms contain all types of observed clusters/ precipitates in one analysis. These proximity histograms are shown in Fig. 5 to a distance of 10 nm from the precipitate center, where the bulk concentration of the Mg crystal is dominating. Towards the center of the precipitates, the xaxis was cut off at 1 nm distance because of poor counting statistics caused by the decreasing number of atoms near the center of the clusters. In Fig. 5a and b, the proximity histograms of the as cast-states of MRI230D and AXZ931 are displayed. It can be seen that the concentrations of Al and Ca increase towards the cluster centers, while no significant increase in Mn seems to be present. What is apparent is that the average concentration of Al is roughly twice that of Ca. This could indicate that these clusters are a precursor of Al2Ca phase. Those prestages, like Guinier Preston zones, were observed in an Mg-0.5Ca-0.3Al (wt.-%) alloy in peak aged condition in Ref. [30] and during creep deformation in an Mg-2.2Al–2Ca-0.3Mn (wt.-%) alloy [31]. Unfortunately, the crystallographic structure of the precipitates could not be assessed using atom probe in our work. This is partially due to local magnification around small precipitates [28], which is also seen to be responsible for the gradual increase in measured concentration towards the cluster centers, rather than a distinct interface as would be expected for a discrete phase. In Fig. 5c and d, which correspond to the crept conditions, a distinctively different situation is encountered. While the ratio of Al:(Ca + Mn) is still close to 2:1, a significant amount of Mn is incorporated in the precipitates. Moreover, significant differences between MRI230D (Fig. 5c) and AXZ931 (Fig. 5d) emerge.While the former predominatley incorporates Mn in the precipitates, the precipiates in the latter contain about equal amounts of Mn and Ca. Also, the overall detected concentrations of solute elements in the clusters is lower in MRI230D, which is an indication of smaller clusters. This is, in tendency, reproduced after ageing treatment (Fig. 5e and f), where for MRI230D (Fig. 5e), Mn dominates in the clusters over Ca. For AXZ931 however, about the same amount of Ca and Mn are incorporated into the precipitates, with slighlty more Ca. As the ageing time is prolonged 4

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between precipitates and bulk is much weaker, measurement artefacts like the artifical detection of two or more precipitates as a single entity is promoted. For the analysis, we used a self-optimizing algorithm based on Voronoi spatial tessellations presented in Ref. [29]. The results of the cluster size analysis are presented in Fig. 6. Here, the total cluster size in solute atoms (Ca + Mn) is plotted versus the atomic fraction of Ca/(Ca + Mn) in each cluster. Additionally, a representative part of the related reconstruction and the cluster density are shown as an inset in each subfigure. To get a better impression of the size of the clusters, an sphere-equivalent-diameter-approximation is given on the second y-axis. This is under the assumption that the precipitates are discrete entities with the cubic C15-Laves structure of the Al2Ca phase, where Ca is substituted by Mn. The lattice parameter of this phase is 0.804 nm [31]. It should be noted that, due to detection efficiency of the atom probe (~37%), very small clusters (< ~11–20 atoms, depending on the measurement) are unaccounted for, as they become undistinguishable from random fluctuations of solute content in the parent crystal. Fig. 6a and b show the cluster analysis of the as-cast alloys of MRI230D and AXZ931, respectively. The clusters consist mainly of Ca and show very similar number densities, with 56 × 10³ clusters/μm³ for AXZ931 and 58 × 10³ clusters/μm³ for MRI230D. The average sphere equivalent diameters of the clusters are ~2.1 and ~2.2 nm respectively. This is also quite similar between the alloys. This changes significantly after creep deformation. After deformation, the clusters tend to be significantly smaller in the case of MRI 230D (Fig. 6c) at an average sphere equivalent diameter of 1.8 nm, while for AXZ931 (Fig. 6d) a broader size distribution with both coarser (around 3 nm) and finer (around 1.85 nm) particles is observed. This implies cutting of the precipitates and/or dissolution and re-formation. Very importantly, Mn gets incorporated in the clusters increasingly with decreasing cluster size. This holds true for both alloys. The broader distribution of cluster sizes in AXZ931 also means that more solute content is tied up in the individual clusters, resulting in a significantly lower number density (82 × 103 μm−3) over the MRI230D alloy (165 × 103 μm−3). Even larger differences between the alloys are present after ageing without mechanical load. This is represented in Fig. 6e for MRI230D and Fig. 6f for AXZ931. While MRI230D shows a precipitate size and chemical distribution similar to the crept condition (Fig. 6c), with a cluster density of 151 × 103 μm−3, significant coarsening is observed for AXZ931, leading to a cluster density of only 40 × 103 μm−3. Most importantly, without mechanical load, the split into coarse, Ca rich precipitates and fine, Mn rich precipitates becomes apparent. This is observed for both MRI230D and AXZ931, but the fraction of coarse precipitates is much larger in the latter. This further underlines the prominent role of Mn in the stabilization of a fine precipitate microstructure. Also, we assert that a change in the strengthening mechanism would occur for AXZ931 if load was applied to the AXZ931 alloy in the aged condition from precipitate cutting to circumventing, because the Ca-rich precipitates have grown quite big compared to as-cast condition. Analysis of the precipitate structure after further ageing to 500 h total ageing time is shown for MRI230D in Fig. 6g. While the duplex structure with fine, Mn rich clusters and coarse Ca rich precipitates is retained, the Mn rich clusters seem to have further decreased in size and increased in numbers (181 × 103 μm−3). This implies that both Ostwald ripening and nucleation of clusters is occurring simultaneously, i.e. the full precipitation potential of the alloy still has not been reached. This is also reflected in the creep curves (Fig. 2), where the minimum creep rate is present at a much later time in MRI230D (~44 h) compared to AXZ931 (~6.5 h). This evolution of the number densities of the clusters are shown in Fig. 7. Here, the number densities of the clusters are plotted against the time the samples remained at elevated temperatures. For MRI230D, the number density steadily increases with time at elevated temperature up

Fig. 5. Concentration calculations (at.-%) (proxigrams) towards the cluster centers of all types of clusters or precipitates observed in one analysis.

to 500 h for MRI230D (Fig. 5g), Mn still shows a larger enrichment in the precipitates than Ca, with Mn to Ca ratios similar to the crept sample but slightly lower than for the 250 h aged sample where Mn seems to be enriched more strongly. While the analysis of the evolution of the average chemical composition already clearly shows some trends and differences between MRI230D and AXZ931, it was already clear from the qualitative, visual analysis of the data presented in Fig. 3 that a spectrum of cluster sizes and possibly compositions exists for the various states of the materials. We therefore employed solute analysis techniques (“cluster search”) in order to quantify the contributions of the alloying elements, specifically Ca and Mn to individual precipitates. The incorporation of Al was not included in the analysis, since it does not show any significant change between the analysed materials. Also, since the partitioning of Al 5

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Fig. 6. Results of the cluster analysis for calcium and manganese atoms. The x-axis shows the relative amount of calcium to the total number of calcium and manganese atoms (Ca/(Ca + Mn))-ratio in each cluster (zero is only manganese atoms and 1 is only calcium atoms). On the left Y-axis the total cluster size in atoms is displayed. The right Y-axis shows a corresponding sphere equivalent diameter of an Al2Ca-precipitate.

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AXZ931 microstructure, there is only a small difference in number density of clusters over the as-cast condition. While a comparable rise in volume fraction up to 2 vol.-% was observed, the number density increased only slightly from 56 × 10³ to 82 × 10³ μm−³ (Fig. 6b and d). However, due to the notable difference in chemistry of the precipitates, it is at this point unclear whether this is the result of simple incorporation of Mn into the existing precipitates or the nucleation of new, Mn containing ones that grow at the expense of Ca rich clusters. Since no “gap” in composition between Mn rich and Ca rich precipitates was found, we interpret this as an evolution where either Mn is gradually incorporated in the precipitate’s crystal lattice or no distinct crystalline second phase is present. The latter would be supported by the fact that no equivalent phase to the Al2Ca phase is present in the Mg–Al–Mn ternary system. If only coherent atom clusters are able to incorporate Mn, which is almost non soluble in the α-Mg matrix (< 0.04 at% at 300 °C) [32,33], Mn would stabilize a fine precipitate structure, which is observed. 4.3. Aged microstructures Fig. 7. Number density of clusters as a function of time at elevated temperatures.

The notion that the incorporation of Mn into the clusters/precipitates stabilizes small, potentially coherent clusters over discrete, incoherent second phases is further supported by an analysis of the microstructure after ageing without stress. In the MRI230D alloy, which contains around 0.17 at.-% Mn, a large amount of the small, Mn rich clusters are retained, while only a few larger Ca rich precipitates are present. These have distinctively separate concentration ranges. In the AXZ931 alloy, with only 0.1 at.-% Mn addition, this stabilization effect is much less pronounced. This leads to the formation of a greater number of large, Ca rich precipitates and with it, a significant decrease in precipitate number density. In comparison to the crept condition (Fig. 3c), the aged MRI230D alloy (Fig. 3e) appears to have a similar microstructure. The volume fraction covered by the isosurfaces in the aged MRI230D is slightly lower at ~1.7 vol.-% (Fig. 4). So is the number density of the clusters (Fig. 6f), being 151 × 10³ clusters/μm³, compared to the crept sample with 165 × 10³ clusters/μm³. In contrast, for the aged AXZ931 sample (Fig. 3f), a complete change in the resulting precipitate structure is observed compared to the crept AXZ931 sample (Fig. 3d). Instead of fine distributed precipitates/clusters like in the crept condition, the aged precipitate structure of AXZ931 contains fewer, relatively large precipitates, as outlined by the 1 at.-% Ca-isosurfaces. Similarly, the Mn-rich areas in AXZ931 changed size and distribution, from being very small and of high number density in crept condition (Fig. 3d and 6d) to larger grown isosurfaces after aging for 250 h (Fig. 3f). Interestingly, in the aged condition, the Mn-rich precipitates appear to form more likely in the vicinity of the larger Ca-rich precipitates. This suggests heterogeneous nucleation at the precipitate/matrix interface and nucleation in precipitate’s strain field. An influence of the precipitate strain field would be supported by the observation that initial nucleation of the Mn rich precipitates is easier, but growth is difficult. Significant lattice strain would cause such a behavior. The volume fraction occupied by the combined Ca and Mn isosurfaces in aged AXZ931 increases to about 2.4 vol.-% (Fig. 4), while the number density is only 40 × 10³ clusters/ μm−³. In order to assess the long term stability of the precipitate microstructure of MRI230D, we included a sample that was aged for 500 h (Fig. 3g). The APT reconstruction of this sample shows a precipiate distribution similar to the MRI230D sample aged for only 250 h, showing that the Mn richer microstructure is more stable over time. However, the Ca rich precipitates, likely to be Al2Ca, show some coarsening, reflected in both the 1% Ca isosurfaces (Fig. 3g) and precipitates found in the cluster search (Fig. 6g). How this coarsening of a small number of Ca rich precipitates influences the stability of the precipitate structure after even longer times is unknown.

to the maximum time investigated in this work, while for AXZ931, differences exist between loaded and unloaded specimens. While for both alloys, the precipitate chemistry changes significantly from the ascast condition, for AXZ931, the number density does not increase as distinct as for MRI230D in crept condition. 4. Discussion 4.1. As-cast condition Starting with the as-cast conditions, it should be noted that using APT, we were able to show the presence of small precipitates in both the AXZ931 and the MRI230D grain interior. This is in contrast to earlier TEM investigations of Amberger et al. [2], where small precipitates were only found in MRI230D. These precipitates, probably early stages of the formation of the Al2Ca phase, seem to not be stable in their initial, almost Mn free form upon exposure to temperatures around 200 °C. Their volume-fraction and number density is also relatively low at around 1 vol.-%, measured by isosurfaces (Fig. 4) and 56 × 10³ clusters/μm³ for AXZ931 and 58 × 10³ clusters/μm³ for MRI230D (see Fig. 6a and b). The corresponding Al/(Ca + Mn)-ratios near the cluster center are ~2.5 for MRI230D and ~2.6 for AXZ931, implying that the clusters are already close to the composition of Al2Ca. 4.2. Microstructures after creep at 200 °C After creep deformation, the precipitate microstructure has significantly shifted through the incorporation of Mn into the particles. This may be explained by the thermal history of the alloys. After casting, the alloys cool from their solidus temperatures, during which precipitation of the Al2Ca phase is possible, while the added amounts of Mn are still soluble in the Mg crystal lattice. Upon long term exposure of the alloys to temperatures around 200 °C, where the solubility of Mn in Mg is close to zero, the Mn that is present in excess in the Mg lattice will seek to precipitate and thus either nucleate new clusters or be incorporated in existing ones. For the MRI230D, creep at 200 °C led to the formation of very fine, Mn rich clusters with a narrow size range, while for AXZ931, both fine and some coarser (~3 nm diameter) precipitates are observed. Beside this, the volume-fraction of the isosurfaces (Fig. 4) increased from ~0.8 vol.-% (as-cast) to more than 2 vol.-% for both alloys. Similarly to the volume-fraction, the number density of the clusters increased from 58 × 10³ to 165 × 10³ μm−³ for MRI230D (see Fig. 6a and c). For the 7

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A positive effect of manganese addition on the creep properties of an ~ Mg–2Al–2Ca (wt.-%) alloy was reported earlier in Ref. [31]. In this work, the improvement in minimum creep rate was reported due to plate like precipitates that were identified by TEM as Al2Ca Lavesphase with a C15-structure that formed during creep experiments. Moreover, the formation of ordered Guinier-Preston-zones (GP-zones) on the basal planes of the α-Mg matrix during creep, which were also identified to contribute to the improvement in creep resistance, were observed. The number densities of the GP zones with 51 × 10−3 μm−3 and 250 × 10−3 μm−3 and an average size of 19.4 nm and 14.3 nm were obtained in the TEM, showing that Mn addition leads to an increase in number density and a decrease of the size of the GP zones. These findings can also be adopted to our results, where the higher Mn content in MRI230D leads to smaller clusters of higher number density, compared to AXZ931. From our results, we also propose that a higher manganese content has a positive influence on the nucleation of new clusters, because the number densities are much lower in AXZ931 (see Fig. 7) having less Mn addition than MRI230D (see Table 1). All this leads us to the conclusion that the small clusters in our alloys act as an effective obstacle for dislocation motion during creep.

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgement The authors would like to acknowledge funding by the Deutsche Forschungsgemeinschaft via the Cluster of Excellence ‘Engineering of Advanced Materials’. References [1] M.K. Kulekci, Magnesium and its alloys applications in automotive industry, Int. J. Adv. Manuf. Technol. 39 (9–10) (Nov. 2008) 851–865. [2] S. Zhu, et al., Evaluation of magnesium die-casting alloys for elevated temperature applications: microstructure, tensile properties, and creep resistance, Metall. Mater. Trans. A 46 (8) (Aug. 2015) 3543–3554. [3] M.M. Avedesian, H. Baker, A.S.M.I.H. Committee, ASM Specialty Handbook: Magnesium and Magnesium Alloys, ASM International, 1999. [4] E. Aghion, B. Bronfin, M. Katzir, S. Schumann, B.F. Von, High Strength Creep Resistant Magnesium Alloy - Patent CA2366610, (03-Jan-2012) CA2366610C. [5] W. Blum, Y.J. Li, X.H. Zeng, P. 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5. Conclusions The precipitate microstructure of the alloys AXZ931 and MRI230D within individual grains were investigated in as-cast, crept and aged conditions via APT. APT measurements revealed the presence of small precipitates in the AXZ931 grain interior, which were not found using TEM imaging in earlier investigations in Ref. [15]. In the as-cast conditions of both investigated alloys, only calciumrich precipitates are present. This is likely owing to the thermal history of the alloys, as Mn solubility in the matrix is high enough for the Mn to be kept in solid solution at elevated temperatures. This changes when application relevant loads and temperatures (200 °C) are applied in the creep tests. After creep, mainly fine, manganese-rich clusters are present in the grain interior. If they comprise a discrete phase with a crystal structure different from Mg or are simply clusters of atoms could not be determined using APT. In these crept states, the biggest differences between MRI230D and AXZ931 are the cluster density, which is higher in MRI230D, and the size distribution, where MRI230D exhibits much smaller clusters. In comparison to the crept sample, the aged conditions of MRI230D still show distributions of very small and manganese-rich clusters, while in AXZ931 significant coarsening sets in. This is attributed to the fact that without mechanical load and the resulting constant production of precipitate nuclei through cutting, Ostwald ripening is significantly accelerated. Still, the crept sample of AXZ931 shows some small manganese-rich clusters with two major differences: (i) The overall cluster density with ~80 × 10³ clusters/μm³ is only half the value compared to MRI230D (165 × 10³ clusters/μm³) and (ii) the cluster size distribution is clearly bimodal with fine Mn-rich and coarser Ca-rich clusters that have an average sphere equivalent diameter of ~1.85 nm and ~3 nm, respectively. Yet even without deformation, the Mn rich clusters in the MRI230D alloy exhibit excellent long term stability (after 500 h aging at 200 °C). The exact mechanistic origin of this effect remains unexplained by the data presented here, but we conclude that these ‘clusters’, rather than precipitates, may not have a distinct crystal structure or a high interfacial energy, which keeps them from coarsening. To facilitate this effect, the Mn concentration present in the grain interior of MRI230D (~0.03 at.-%) seems to be sufficient. How this effect changes with the Mn concentration remains to be investigated.

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Sci. 246 (1) (Apr. 1991) 442–449. [29] P. Felfer, A.V. Ceguerra, S.P. Ringer, J.M. Cairney, Detecting and extracting clusters in atom probe data: a simple, automated method using Voronoi cells, Ultramicroscopy 150 (Mar. 2015) 30–36. [30] J. Jayaraj, C.L. Mendis, T. Ohkubo, K. Oh-ishi, K. Hono, Enhanced precipitation hardening of Mg–Ca alloy by Al addition, Scripta Mater. 63 (8) (Oct. 2010) 831–834.

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