Materials Science & Engineering A 558 (2012) 403–411

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Materials Science & Engineering A

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journal homepage: www.elsevier.com/locate/msea

Factors influencing creep flow and ductility in ultrafine-grained metals

V. Sklenicka a,b,n, J. Dvorak a,b, P. Kral a, M. Svoboda a,b, M. Kvapilova a, T.G. Langdon c,d

a Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, CZ-61662 Brno, Czech Republicb CEITEC—IPM, Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, CZ-616 62 Brno, Czech Republicc Departments of Aerospace & Mechanical Engineering and Materials Science, University of Southern California, Los Angeles, CA 90089-1453, USAd Materials Research Group, Faculty of Engineering and the Environment, University of Southampton, Southampton SO17 1BJ, UK

a r t i c l e i n f o

Article history:

Received 3 July 2012

Received in revised form

3 August 2012

Accepted 6 August 2012Available online 10 August 2012

Keywords:

Creep

Ductility

Equal-channel angular pressing

Microstructure

Ultrafine-grained material

93/$ - see front matter & 2012 Elsevier B.V. A

x.doi.org/10.1016/j.msea.2012.08.019

esponding author at: Institute of Physics of M

zech Republic, Zizkova 22, CZ-61662 Brno, Cz

0; fax: þ420 541 212 301.

ail addresses: sklen@ipm.cz, vsklen@ipm.cz (V

ipm.cz (J. Dvorak), pkral@ipm.cz (P. Kral), sv

@soton.ac.uk, langdon@usc.edu (T.G. Langdo

a b s t r a c t

The creep behaviour of high purity aluminium and copper and their Al–0.2 wt%Sc and Cu–0.2 wt%Zr

alloys was examined after processing by equal-channel angular pressing (ECAP) with an emphasis on

creep ductility and the ECAP microstructural homogeneity. It was found that, under the same loading

conditions, the creep ductility of the ultrafine-grained materials processed by ECAP continually

increases with increasing numbers of ECAP passes. A detailed quantitative microstructural study was

conducted using the electron backscatter diffraction (EBSD) methods. This analysis revealed that, with

increasing numbers of ECAP passes, the mutual misorientation of neighbouring subgrains grows and

the subgrains continuously transform to grains having high-angle grain boundaries.

& 2012 Elsevier B.V. All rights reserved.

1. Introduction

Creep ductility is very important for various shaping and formingtechnological operations at elevated and high temperatures andespecially for avoiding catastrophic failure in the load-bearing partsof high temperature components. Creep strength and ductility arethe key creep properties of creep-resistant materials but theseproperties typically have opposing characteristics. Thus, thesematerials may be strong or ductile but they are rarely both. In thisconnection, recent findings of high strength and good ductility inseveral bulk ultrafine-grained (UFG) metals produced by severeplastic deformation (SPD) are of special interest [1].

Several (SPD) processing techniques [2–5] are currently avail-able but the most attractive technique is equal-channel angularpressing (ECAP) where the billet is pressed through a die con-strained within a channel bent though an abrupt angle [1].Processing by ECAP provides an opportunity to achieve verysignificant grain refinement to the submicrometer or even thenanometer level. Thus, there is a potential for using pressedmaterials to obtain new flow processes in high temperature creepprovided the ultrafine grains are reasonably stable at elevatedtemperatures. At present only very limited reports are available

ll rights reserved.

aterials, Academy of Sciences

ech Republic. Tel.: þ420 541

. Sklenicka),

obm@ipm.cz (M. Svoboda),

n).

describing the creep behaviour of metals and alloys processed byECAP and then tested under conditions of constant stress orconstant load [6–11] and most of these investigations wereinitiated to provide basic information on the rate-controllingcreep mechanism(s) and the relevant steady-state and/or mini-mum creep rates. However, neither the phenomenological nor themicroscopic aspects of creep ductility of UFG materials areunderstood at the present time.

A strong influence of the microstructure on creep behaviourwas observed in various UFG materials [7,8,12,13]. By contrast, noreport is available describing the link between microstructure andcreep ductility in UFG materials processed by ECAP at elevatedand high temperatures. Thus, the aim of the present work was toobtain a better insight into the effect of the microstructureproduced by ECAP on the subsequent creep ductility of alumi-nium, copper and some of their precipitation-strengthened alloys.The first section summarizes the creep properties of thesematerials and the following section describes results from aqualitative and quantitative microstructural examination. Themicrostructure was revealed by electron backscatter diffraction(EBSD) and analysed by stereological methods. The effect ofrepeated ECAP passes on creep behaviour and microstructuralvariables is examined in detail.

2. Experimental materials and procedures

Earlier reports described the creep properties of high-purity(99.99%) aluminium processed by ECAP [6,14]. In those

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411404

investigations, the initial extremely coarse grain size of �5 mmwas reduced to �0.8 mm by pressing for 8 passes at roomtemperature using an ECAP die containing a channel bent throughan angle of 901 and with an outer arc of curvature of 201 at thepoint of intersection of the two channels. It can be shown thatthese angles lead to an imposed strain of �1 on each separatepassage through the die [15]. The samples were processed usingroute BC in which each billet is rotated by 901 in the same senseabout the longitudinal axis between subsequent passes [16]. Inpractice, the processing route BC was used because it leads mosteffectively to an array of equiaxed grains separated by bound-aries having high angles of misorientation [17]. All other experi-mental materials investigated in this work were processed atroom temperature using the same ECAP die and processing routeBC. The high-purity (99.99%) copper was received in a coarse-grained state with a grain size of �1.2 mm. Repetitive ECAPpressing was conducted to give totals of up to 12 passes. Thegrain size of pure copper after four passes using route BC wasmeasured as �0.4 mm [10].

It should be emphasized that the ultrafine grain sizes of thesetwo pure metals produced by SPD processing are not stable whentesting in high temperature creep [10,14]. This was the reason foralso using two precipitation-hardened Al–Sc and Cu–Zr alloys.A scandium addition of �0.2 wt%Sc to pure aluminium and azirconium addition of �0.2 wt%Zr to pure copper are sufficient toreasonably retain a small grain size at selected creep testingtemperatures [18]. An Al–0.2 wt%Sc alloy with an initial

Fig. 1. Creep curves of pure aluminium for unpressed state and various number of ECAP

(b) creep rate de/dt vs. normalized time t/tf, and (c) creep strain e vs. time t/tf.

extremely coarse grain size of 5–10 mm was processed by ECAPfor 8 passes to give a grain size of �0.55 mm. Full details on thefabrication of this alloy are given elsewhere [19,20]. A coarse-grained Cu–0.2 wt%Zr alloy was homogenized for 24 h at 1073 Kand hot-rolled. Before ECAP, the billets were solution treated at1233 K for 1 h to give an initial grain size of �350 mm. Thepressing was conducted at room temperature up to 12 ECAPpasses to give a grain size of �0.35 mm [21].

Constant load creep tests were conducted in tension usingspecimens having gauge lengths of 10 mm and cross-sectionalareas of 8�3.2 mm2. Each creep specimen was machined so thatits longitudinal axis was oriented parallel to the pressing axis. Thecreep testing was conducted in an environment of purified argonwith the testing temperatures maintained to within 70.5 K of thedesired value. All of the tests were continued until final fracture.For comparison purposes, additional creep tests were conductedon specimens of the same but unpressed (coarse-grained) materi-als. The strain-time readings were continuously recorded using aPC-based data acquisition system.

The application of modern imaging methods to the quantitativeexamination of microstructures in UFG materials processed byECAP permits a more detailed investigation of a possible linkbetween the internal microstructures and the mechanical beha-viour [1,12]. Diffraction-based techniques for localized crystalorientation measurements, such as EBSD, are of central importancetoday for a characterization of these fine-scale microstructuralfeatures [22–24].

passes (creep in tension up to fracture): (a) standard creep curves (data from [6]),

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411 405

In addition to the grain size determination, there are a numberof important microstructural parameters evaluated from EBSDbut not available from conventional methods of grain character-ization, in particular parameters relating to the grain orientationsand boundary characters [25].

For each examined specimen, three mutually perpendicularplanar metallographic sections were selected denoted as XY, XZ

(longitudinal sections) and YZ (transverse section), where X, Y andZ are the axes of the Cartesian coordinate system with X along thelast pressing direction and Z perpendicular to the bottom of thechannel. The technique of automated EBSD in the scanningelectron microscope was used for quantitative metallography.Four ranges of the boundary misorientation D were selected;21rD, 51rD, 101rD and 151rD. Then standard interceptcounting was used [26] to give the mean number NL of profilechords per unit length of the examined test lines. In each speci-men, six systematically selected directions of the test lines in eachsection were examined. The mean boundary areas unit volumewere estimated by the stereological relation SV¼[2NL].

Another important feature of the grain boundary structure isits inhomogeneity. The dispersion of grain profile areas can bequalified by the coefficient of variation CVa of the grain profileareas in a plane CVa ¼

ffiffiffiffi

Vp

=x, where V is the grain profile areasvariation and x is the mean value of the grain profile area [26–28].The coefficient of variation CVa of the profile areas is perhaps thebest stereometric characteristic to evaluate homogeneity ofmicrostructure and today it is relatively easily attainable usingcomputer image analysis.

Fig. 2. Creep curves of pure copper for unpressed state and various number of ECAP pa

de/dt vs. normalized time t/tf, and (c) creep strain e vs. time t/tf.

3. Experimental results

3.1. Creep behaviour

Representative standard creep strain e versus time t curves areshown in Figs. 1a–4a: these plots were obtained at a temperatureof 473 K (�0.5Tm) for Al and the Al–0.2Sc alloy, 473 K (�0.35Tm)for Cu and 673 K (�0.5Tm) for the Cu–0.2Zr alloy. For eachmaterial the specimens with different numbers of ECAP passeswere loaded under the same initial applied uniaxial tensile stressand the creep tests were run up to the final fracture of the creepspecimens. It is important to note that there is a difference in theappearance of the creep curves between the unpressed and thepressed materials and there is a difference in the fracture strainlevels for the pressed material with different numbers of ECAPpasses: these differences are denoted by the numbers B1–B12where the numeral denotes the number of ECAP passes.

Several important conclusions may be reached from Figs. 1a–4a. The unpressed materials exhibit markedly different creep lifeor markedly different duration of creep exposure than thematerials processed by ECAP. The same applies for specimens ofthe same material with different numbers of ECAP passes wherethe samples labelled B1 consistently show the lowest creep rates.To investigate the creep strain during creep exposure and formutual comparison of the fracture strains, Figs. 1a–4a arereplotted in the form of the instantaneous strain rate de/dt and/orthe creep strain e versus the creep exposure time t normalized to thetime to fracture tf as shown in Figs. 1b, c–4b, c. As demonstrated

sses (creep in tension up to fracture): (a) standard creep curves [10], (b) creep rate

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411406

by Fig. 1b and 2b, the minimum creep rates for the pressedpure aluminium and copper are about one to two orders ofmagnitude lower than for their counterparts depending on thenumbers of ECAP passes. The situation with precipitation-hardened Al–0.2Sc and Cu–0.2Zr alloys are more complicated asshown in Figs. 3b and 4b. The minimum creep rate of theunpressed Al–0.2Sc alloy is comparable with that for the pressedalloy after the first two ECAP passes whereas for the pressedCu–0.2Zr alloy the minimum creep rates after one and two ECAPpasses are even slower than the minimum creep rate correspond-ing to the unpressed state. Nevertheless, at higher numbers ofECAP passes the minimum creep rates of both pressed alloysbecome faster than those for their unpressed counterparts andthese differences consistently increase with increasing numbersof ECAP passes.

All materials generally exhibit three distinct regions of flow.The primary region corresponding to the initial stage of creepwhere the creep rate decreases due to work hardening, thesecondary or steady stage (which is frequently reduced to aninflection point in the de/dt versus t/tf curve) in which the creeprate remains approximately constant and the tertiary stage wherethe creep rate accelerates to final fracture. There are furtherdifferences in the creep behaviour of these materials in theaccumulation of the creep strain during the course of creepexposures as shown in Figs. 1c–4c. The figures demonstrate thatvery significant strain contributions to the fracture strain aregenerated during the last 10% of the creep life.

Fig. 3. Creep curves of an Al–0.2 wt%Sc alloy for unpressed state and various number

(b) creep rate de/dt vs. normalized time t/tf, and (c) creep strain e vs. time t/tf.

3.2. Microstructural investigations

There are numerous reports of the processing of various puremetals and metallic alloys by ECAP and many of these reportsinvolve a detailed characterization of the microstructure. Theseresults are summarized in recent reviews [1,2]. However, infor-mation is seldom reported on the percentage of high-angle grainboundaries (HAGBs) although this is an important parameter inany comparison of plasticity of different processing routes andmaterials [29]. It can be expected that samples with differentdistributions of misorientations across the grain boundaries andtextures will deform differently [30]. Furthermore, to provideinformation on the optimum microstructure of UFG materials, itis necessary to use an additional quantitative microstructuralparameter other than the average critical grain size for theenhancement in creep ductility [31]. Such parameter could bethe coefficient of variation CVa of the profile areas as a measure ofthe homogeneity of the material microstructure [28].

Accordingly, the grain and subgrain structures of the creepspecimens were revealed by EBSD and characterized by thecoefficient of variation CVa of the profile areas. Four ranges ofthe boundary misorientation D between adjacent pixels wereselected for examination using EBSD, which correspond specifi-cally to subboundaries, transitive and high-angle grain bound-aries within 21rD and 51rD, transitive and high-angle grainboundaries for 101rD, and mostly grains with HAGB�s for DZ151.Selected examples of images of XZ sections produced by EBSD of

of ECAP passes (creep in tension up to fracture): (a) standard creep curves [19],

Fig. 4. Creep curves of a Cu–0.2 wt%Zr alloy for unpressed state and various number of ECAP passes (creep in tension up to fracture): (a) standard creep curves [21],

(b) creep rate de/dt vs. normalized time t/tf, and (c) creep strain e vs. time t/tf.

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411 407

an Al–0.2 wt%Sc are shown in Fig. 5. It was generally observedthat, with increasing numbers of ECAP passes N, a considerablefraction of subgrain boundaries is gradually transformed to high-angle grain boundaries as shown in Fig. 6. At the same time, thelocal homogeneity of structure as characterized by the values ofCVa with the increasing number N is gradually improved. This isdemonstrated in Fig. 7 which shows the microscopic appearanceof specimens of an Al–0.2 wt%Sc alloy crept under the sameloading conditions (473 K, 20 MPa) but processed by ECAP withdifferent numbers of ECAP passes. Values of CVa as high as 10 atN¼2, 1rCVao2 at N¼4, and 0.55rCVar1 at N¼8 were found.The extremely high value of CVa at N¼2 demonstrates the veryhigh inhomogeneity of a mixture of subgrain and grain structuressince the value of CVa in a very homogeneous grain system shouldnot exceed a value of 1 [12]. In such circumstances, a greatvariability of creep fracture ductility is a natural consequence ofthe short as well as long-range inhomogeneity of the microstruc-ture of individual tested specimens.

A substantial grain coarsening developed, especially in thecase of pure metals, during the creep exposures depending on thestress and temperature, thus manifesting the thermal instabilityof the ultrafine-grained microstructure [1,18,32]. It is clear that inpure aluminium and copper the grains grow rapidly at elevatedtemperatures because there are no precipitates within the crys-talline lattice to restrict the movement of the grain boundaries bya pinning effect. By contrast, submicrometer grains may beretained to relatively high temperatures in materials containing

a distribution of fine precipitates; thus, an Al–0.2%Sc alloy con-tains Al3Sc precipitates [33] and a Cu–0.2%Zr alloy containsCu9Zr2 precipitates [21]. However, from the qualitative point ofview for the subgrain and grain shapes and profile dispersion, thestructures developed under different creep loading conditions ofthese alloys are similar [13].

A quantitative characterization of the inhomogeneity of theboundary structure is shown in Fig. 8. It should be emphasizedthat the values of the coefficient of variation of the profile areaCVa as a measure of structure homogeneity strongly depend onthe chosen ranges of misorientation D in the EBSD analysis.Whereas the fraction of the subboundaries (low-angle grainboundaries) are dominating for DZ21 (Fig. 9a), the fractions ofhigh-angle grain boundaries (yZ151) confirm their high share forthe range of DZ151 (Fig. 9b). Detailed inspection of Fig. 8b showsa systematic dependence of CVa after creep on the numbers ofECAP passes. A substantial decrease of the values CVa 21 for bothprecipitation-hardened Al–0.2Sc and Cu–0.2Zr alloys withincreasing number of passes for DZ21 may be connected withthe more rapid evolution of boundaries having misorientationangles y4151 (Fig. 8a).

4. Discussion

A knowledge of the operating creep deformation mechanismsin UFG materials assists in understanding creep ductility in these

Fig. 5. Selected examples of EBSD maps of an Al–0.2 wt%Sc alloy after ECAP taken

at different ranges of EBSD misorientation D showing: (a) subboundaries DZ21,

(b) transitive subboundaries and high-angle boundaries DZ101 and, (c) mostly

high-angle grain boundaries DZ151.

Fig. 6. The fracture of high-angle grain boundaries in the crept samples as a

function of the number of ECAP passes.

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411408

materials. The total creep strain generally consists of severalcontributions:

e¼ egþegbsþedþev, ð1Þ

where eg, egbs, ed, ev are the strains caused by intragranulardislocation processes including dislocation glide and climb, grainboundary sliding (GBS), stress directed diffusion of vacancies andby intergranular void nucleation and growth, respectively. Itshould be noted that not all of the processes operating areindependent of each other. Recent work has analysed the flowmechanisms occurring in ultrafine-grained aluminium, copperand their selected alloys [13,34–36]. It was found that the

ultrafine grains produced by ECAP in high purity aluminium andcopper are not sufficiently stable when testing in high tempera-ture creep so that the dominant creep process is an intragranulardislocation mechanism as in metals with coarse grain sizes. Anincrease in the fraction of HAGB’s with increasing number of ECAPpasses, as in Fig. 6, will essentially lead to an increasing contribu-tion of GBS to the total creep strain as was determined experi-mentally [6]. Nevertheless, it was found that for pure aluminiumprocessed by 12 ECAP passes and crept at 473 K and 15 MPa, thehighest contribution of GBS to the creep fracture strain was onlyabout 30%. To avoid the occurrence of significant grain growthand thermal instability in creep tests conducted on high-purityaluminium and copper after processing by ECAP, further analyseswere undertaken to examine the flow characteristics ofprecipitation-hardened aluminium and copper alloys where scan-dium and/or zirconium were added to retain the ultrafine grainsizes at elevated temperatures. It was confirmed that scandiumadditions of 0.2 wt% to pure Al are sufficient to retain a smallgrain size at elevated temperatures and this leads to the conclu-sion that flow occurs mostly by GBS as in superplastic deforma-tion [34].

Preliminary results indicate that the ultrafine grains in Cu–0.2 wt%Zr crept at 673 K are more stable at elevated temperaturesthan Al–0.2 wt%Sc crept at 473 K but nevertheless the precipitatesin this alloy exhibit an effective pinning effect on the grainboundaries which restricts grain growth during creep exposure.Although an intragranular dislocation mechanism appears to bedominating in the Cu–0.2 wt%Zr alloy under the creep conditionsinvestigated, the creep behaviour of this alloy may be influencedby a synergistic effect of additional creep mechanisms like grainboundary sliding and intergranular cavitation [13]. Althoughthere have been a few studies reporting the importance ofdiffusional creep in ultrafine-grained metals [8,37–39], the ana-lyses of the stress dependences of the minimum creep rates ofUFG aluminium alloys [34] and pure copper [36] indicate thatdiffusional creep is not able to account for the creep behaviour inthese materials. In fact, these analyses have shown that experi-mentally determined creep rates are much faster than thosecalculated from the diffusion creep equations.

The characteristics of microstructural evolution in pure Alduring ECAP processing are now well documented [40–42] andthis information may be used to draw conclusions concerning the

Fig. 7. Microscopic appearance of an Al–0.2 wt%Sc alloy processed by different

number N of ECAP passes and crept at 473 K and 20 MPa: (a) 2 ECAP passes,

CVa442, (b) 4 ECAP, CVao2 and (c) 8 ECAP passes, 0.55rCVao1.

Fig. 8. Coefficient of profile area CVa as a measure of homogeneity: 0.55rCVao1

(homogeneous system), and CVa442 (multimodal grain size distribution). The

chosen ranges of misorientation D in EBSD analysis: (a) D¼21 and (b) D¼151.

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411 409

creep behaviour. After the first ECAP pass, low-angle boundariesare predominantly formed as shown in Fig. 6. However, it is alsoknown from studies of materials processed by ECAP that thestructure at this early stage consists of highly-deformed subgrainscontaining very high densities of dislocations [43–45]. Thesehighly-deformed structures inhibit the occurrence of intragranu-lar dislocation processes and this leads to very slow creep rates asdocumented in Figs. 1b–4b and, consequently, to long creep lives.During subsequent passes, the average boundary misorientationsincrease, the structure within the grains becomes more uniformand also the creep rates increase but with a correspondingdecrease in the overall creep life. It is apparent also that theoverall ductility as represented by the strain to fracture exhibits

differing trends. There is an additional increase in ductility withincreasing numbers of ECAP passes and therefore with theimposition of even larger strains. Thus, materials processed byECAP may be more creep-resistant or ductile but they are notsimultaneously both.

It appears that the ductility enhancement after higher num-bers of ECAP passes is associated with the increase in the fractionof high-angle grain boundaries and with a consequent change inthe flow process due to the increasing significance of grainboundary sliding. It is now well established that grain boundarysliding occurs in UFG materials due to the presence of an excess ofextrinsic dislocations in the non-equilibrium grain boundariesand this effect has been demonstrated using both macroscopicUFG materials [46,47] and through the deformation of micro-pillars using nanocompression in a scanning electron microscope[48]. Fig. 10 shows the measured strain to fracture as a function ofthe fraction of HAGBs when creep testing under the creep loadingconditions presented in Fig. 6 and it is readily apparent thatspecimens with higher fractions of high-angle grain boundariesexhibit higher ductility.

Finally, it is necessary to insert a word of caution. Due to thefrequent use of miniaturized tensile specimens in research on

Fig. 9. Distribution of boundaries with different misorientation y for an Al–

0.2 wt%Sc alloy analysed in Fig. 8: (a) EBSD analysis for D¼21 and (b) D¼151.

Fig. 10. Dependence of the strain to fracture ef on the fraction of high-angle grain

boundaries y4151.

V. Sklenicka et al. / Materials Science & Engineering A 558 (2012) 403–411410

UFG materials, the specimen dimensions and/or geometry effectsmay have a significant effect on the measured mechanical proper-ties and overall ductilities [49]. Therefore, although the creepresults obtained in this investigation are mutually consistent, caremust be taken in extending these results to other UFG materialswhere the creep ductilities may differ due to the use of differentspecimen geometries.

5. Summary and conclusions

1.

High purity (4 N) aluminium and copper and their Al–0.2%Scand Cu–0.2%Zr alloys were processed by equal-channel angu-lar pressing (ECAP) and then examined in terms of their creepductility and microstructural homogeneity.

2.

The results show that both the creep rates and the strains tofracture of these ultrafine-grained materials increase withincreasing numbers of ECAP passes. The low creep rates inthe early stages of ECAP processing are due to the very highdislocation densities that are introduced during the pressingoperation.

3.

Microstructural investigations showed that a homogeneousmicrostructure together with a higher fraction of high-anglegrain boundaries leads to an increase in the creep ductility.This is attributed to the increasing significance of grainboundary sliding.

4.

The coefficient CVa of the profile areas is probably the bestquantitative strereometric parameter to characterize thehomogeneity of the microstructure. It is relatively easilyattainable by computer image analysis.

Acknowledgements

The authors acknowledge financial support for this work providedby the Czech Science Foundation under Grant no. P108/11/2260. Thiswork was realized in CEITEC—Central European Institute of Technol-ogy with research infrastructure supported by the project CZ.1.05/1.1.00/02.0068 financed from the European Regional DevelopmentFund. One of the authors (TGL) acknowledges support from theEuropean Research Council under ERC Grant Agreement no. 267464-SPDMETALS.

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