1. Introduction
Since vehicle mass is a key contributor to fuel efficiency, the use of lightweight materials can have a significant impact on their CO
2 emissions. The increased use of magnesium alloys is therefore very appealing, as magnesium is the least dense metal regularly used in structural applications. However, the most widely used magnesium alloys are based on the Mg–Al system and their application is limited to temperatures below approximately 150 °C as a result of the low creep strength of these alloys [
1,
2,
3]. Conversely, e.g., for automotive powertrain applications, alloys with an increased creep strength for use up to 200 °C under long-term loading conditions are needed. Several successful attempts have been made to increase the creep strength of Mg–Al alloys by adding rare earth elements or alkaline earth elements, see for example [
4,
5,
6,
7,
8]. However, there is a lack of understanding and an ongoing debate in the literature on the mechanisms responsible for the increased creep strength.
The microstructure of all derivatives of the above-mentioned Mg–Al-alloy family shows a skeleton-like structure of hard intermetallic phases (IP) in which the Mg grains are embedded. In several publications, it is proposed that the skeleton of hard IPs is responsible for the improved creep strength of these alloys, see [
9,
10,
11,
12,
13,
14,
15,
16]. For Ca-alloyed A91 variants it could be shown that an increased interconnectivity of the IP network results in an increased creep strength [
17]. This is supported by the works of Zubair et al. [
18,
19] where the skeleton of hard IPs was claimed to be responsible for the improved creep properties of the investigated Mg–Al-Ca alloys. Their microstructural investigations of creep-deformed specimens clearly showed strain localizations at the interfaces of α-Mg and IPs and a fracture of the IP in areas of high strain localization [
18,
19].
However, a comparison of the Ca-alloyed AZ91 variants [
17,
20] with the commercial alloy MRI 230D shows that the alloy MRI 230D has an improved creep strength over similar Ca-alloyed AZ91 variants. Even an AZ91 alloy with 5 wt.% Ca addition (AXZ951) has a lower creep strength compared to MRI 230D, especially in a regime of low creep rates [
17]. This is quite surprising, considering MRI 230D has a lower content of Ca, only about 2 wt.%, compared to approximately 5 wt.% in the case of the alloy AXZ951, i.e., a smaller amount of IP. The quantification of the interconnectivity of the IP skeleton additionally revealed that the interconnectivity is less pronounced than in the case of the alloy AXZ951; the interconnectivity of the IP of the alloy MRI 230D is more equal to that of the alloy AXZ931 with 3 wt.% of Ca—see [
17]. Since the creep rate of the alloy AXZ931 is more than one order of magnitude higher than the creep rate of the alloy MRI 230D under identical conditions [
17], the question about the origin of the improved creep strength of MRI 230D arises. Recently, the increased creep strength of the Mg grains due to precipitation was investigated by Lamm et al. [
21]. They showed that MRI 230D has much more thermally stable nano-sized precipitates in the Mg grains, at least partially explaining the increased creep resistance. What is however still unclear is if the IP network also behaves differently in MRI 230D vs. the AZ91 variants.
Light microscopy images of the microstructure of the investigated alloys are summarized in
Figure 1. The microstructure of all alloys consists of α magnesium grains (light grey) surrounded by intermetallic phases (dark grey). The IP shows an increasing interconnectivity with increasing Ca content. Due to the lower thixomolding temperature of the alloy MRI 230D, primary solid α-Mg can be found in the microstructure. A detailed analysis of the microstructure of the investigated alloys including the reproducibility of the microstructural observations as well as the quantification of the interconnectivity can be found in [
17].
As known from previous work, the creep resistance of AZ91 alloy variants increases with increasing Ca content; see for example [
1,
17,
20,
22,
23,
24,
25]. The alloys tested in this work show the same behaviour, as depicted in
Figure 2, where their respective creep curves are shown. The observed strain rate ϵ decreases significantly with an increasing Ca content. An increase in the Ca content from 1 wt.% to 3 wt.% leads to a decrease in the measured minimum creep rate of about one order of magnitude. A further increase in the Ca content to approximately 5 wt.% further decreases the minimum creep rate by more than one order of magnitude. The creep rate of MRI 230D is comparable to that of AZ91 with 5 wt.% Ca addition, although the Ca content of MRI 230D is much lower with only about 2 wt.%; see also [
17]. As the ductility of AZ91 decreases with increasing Ca content (see [
22]) the tensile elongation of the alloy with 5 wt.% of Ca is below 1%. A lower Ca content of the alloys is therefore beneficial for structural applications.
Other works show that a combined addition of Ca and Sr to Mg–Al alloys leads to improved creep properties when compared with the addition of only Sr or Ca [
26,
27], but no reasons for these improved properties due to a combined addition of Sr and Ca are given.
To gain more insight into the mechanisms responsible for the enhanced creep resistance, in this work the mechanical properties of the individual phases have been determined via nanoindentation. To this end, a commercially sourced sample of MRI 230D has been investigated in different states of creep deformation, namely as-cast, at minimum creep rate and after long-term creep loading at a typical service temperature of 200 °C (approximately 250 h). These results are compared with a series of experimental AZ91 alloy variants with Ca additions varying from 1–5 wt.%. The nanoindentation measurements have been performed at room temperature and at 200 °C. Measuring the local mechanical properties of the individual phases is a new approach to gain more insight into the origins of the creep resistance of magnesium alloys, especially in order to understand the high creep strength of the alloy MRI 230D.
In the case of the Ca-alloyed AZ91 variants, the IP mainly consists of Mg
17Al
12, Al
2Ca and (Mg, Al)
2Ca; see for example [
28,
29,
30]. X-ray diffraction measurements from Aghion et al. [
31] revealed only the presence of α-Mg and Al
2Ca in the case of the alloy MRI 230D. Other groups report ternary MgAlSrCa phase compositions in Mg-Al-Ca-Sr alloys; see for example [
32,
33,
34,
35].
2. Experimental Section
A commercially available alloy, MRI 230D, was investigated and compared with Ca-alloyed AZ91 alloys with nominally 1, 3, and 5 wt.% of added Ca, named AXZ911, AXZ931 and AXZ951, respectively, according to ASTM B275 standard. All alloys were thixomolded by Neue Materialien Fürth GmbH on a Japan Steel Works machine with a closing force of 220 t. The cooling rate in thixomolding is rather high, comparable with the cooling rate of high-pressure die casting, and should be comparable for all alloys investigated in this work. In the case of the Ca-alloyed AZ91 variants, a thixomolding temperature of 605 °C was used, so these alloys were cast in a fully liquid state. For the alloy MRI 230D the thixomolding temperature was 590 °C, so the alloy was cast in semi-solid state Due to this, primary α-Mg can be found in the microstructure; see also
Figure 1. Comparative measurements showed no significant influence of the thixomolding temperature on the hardness of the individual phases. The chemical composition for all investigated alloys as measured by GDOES (Glow Discharge Optical Emission Spectroscopy) is summarized in
Table 1.
Throughout this work, great care has been taken to take the samples out of the centre of the casting plates and to avoid taking samples out of the regions of the casting skin.
Creep tests were performed in compression mode at a temperature of 200 °C under constant true stress of 100 MPa on samples of approximately 5.5 mm height and about 20 mm2 in cross-section. The cylindrical samples were ground to a parallelism of better than 10 µm. The solid cross-section, where casting porosity has been taken into account, is calculated from the sample mass m and the sample height h0,RT (averaged over five measurements) as A0,RT = m/(ρh0,RT). A constant solid density ρ = 1.811 g/cm3 was assumed for all alloys. For the calculation of the sample cross-section at the elevated test temperature of 200 °C the thermal expansion was accounted for with a thermal expansion coefficient of 2.6 × 10−5 K−1.
The compressive constant true stress experiments were carried out using a lever-arm creep-testing machine, with the height change Δh measured by a tube–rod extensometer system based on a linear variable differential transformer. Compressive strain is determined as ε = −ln(1 + Δh/h0). Compressive stress is calculated as σ = −F/A0 × exp(−ε) from the force, F, measured by a load cell with a maximum capacity of 20 kN. Please note that for reasons of clarity in this work for compressive strains and stresses absolute values are used. For all creep tests in this work a temperature of 200 °C and a stress of 100 MPa has been used.
For microstructural investigations and nanoindentation measurements, samples were cut out of the castings or the creep-deformed samples and ground and polished with diamond suspension. The last preparation step was polishing with colloidal SiO2 suspension (OPS, Struers, Denmark) to ensure high surface quality with low residual deformation. To prevent corrosion, water-free isopropyl alcohol was used as a lubricant for all polishing steps. In the case of the creep deformed samples, a surface parallel to the deformation direction was investigated.
The nanoindentation measurements at room temperature (meaning a temperature range of 18 °C–25 °C) were performed on an Agilent Nanoindenter XP. For the measurements at 200 °C a G200 from Surface equipped with a heating stage was used. To prevent the oxidation of the samples, measurements at 200 °C were performed under a protective argon atmosphere. All nanoindentation experiments used a Berkovich indenter tip and continuous stiffness mode (CSM). Due to the limited size of the intermetallic phase, the maximum indentation depth was 250 nm and the hardness for each measurement was averaged over a depth range between 200 and 250 nm.
Figure 2 shows exemplarily indents in a sample of the alloy AXZ911 which has crept to the minimum creep rate.
Atom probe tomography (APT) experiments were carried out in a Cameca LEAP 4000X HR (Cameca Instruments Inc., Madison, WI, USA) with a reflectron detector system offering 37% detection efficiency. The samples were prepared via FIB lift-out in a FEI Helios 660i (FEI, Hillsboro, FL, USA) and a Zeiss Crossbeam 1540 EsB (Carl Zeiss AG, Oberkochen, Germany). 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 channelling. For the final milling step, the ion energy was reduced to 5 kV. A more precise description of the lift-out procedures is given in Refs. [
36,
37,
38]. All APT measurements were performed using high voltage pulses with a repetition frequency of 200 kHz to trigger field evaporation with the sample bias voltage regulated to achieve an average detection rate of 2 kHz. The temperature of the tips was 40–45 K and the pulse amplitude was 20% of the bias voltage.
Author Contributions
Conceptualization, S.L., H.W.H. and P.F.; methodology, S.L.; H.W.H. and P.F., validation, D.M.-A., H.W.H., M.G. and P.F.; investigation, P.T. and S.L.; resources, M.G. and P.F.; data curation, D.M.-A., S.L. and P.T.; writing—original draft preparation, D.M.-A.; writing—review and editing, H.W.H., P.F. and M.G.; visualization, D.M.-A., S.L. and P.T.; supervision, H.W.H., P.F. and M.G.; project administration, H.W.H. and P.F.; funding acquisition, M.G. and P.F. All authors have read and agreed to the published version of the manuscript.
Funding
This research was funded by the Deutsche Forschungsgemeinschaft (DFG) via the Cluster of Excellence EXC315 ‘Engineering of Advanced Materials’.
Institutional Review Board Statement
Not applicable.
Informed Consent Statement
Not applicable.
Data Availability Statement
The data presented in this study are available on request.
Acknowledgments
The authors want to thank the Fraunhofer IISB for access to the G200 and the possibility to conduct the nanoindentation measurements at 200 °C.
Conflicts of Interest
The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.
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