8
Segregation and clustering of solutes at grain boundaries in Mg–rare earth solid solutions M. Bugnet a , A. Kula a,b , M. Niewczas a,, G.A. Botton a,a Department of Materials Science and Engineering, McMaster University, 1280 Main Street West, Hamilton, Ontario L8S 4L7, Canada b AGH-University of Science and Technology, Cracow, Poland Received 30 January 2014; received in revised form 27 May 2014; accepted 2 June 2014 Abstract The present study validates the previously reported investigations about segregation of rare-earth (RE) elements at grain boundaries in Mg–RE alloys and ultimately provides a direct visualization of the distribution of the solute atoms in the structure of a Mg–Gd alloy. It is demonstrated that Gd forms a solid solution within the Mg matrix in addition to substantial segregation at high-angle grain bound- aries in the form of 1–2 nm clusters, with a postulated face-centered cubic Gd structure. The results suggest significant implications for the texture development during alloy processing and recrystallization, and thus for the mechanical behavior and properties of Mg–RE alloys. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium–rare earth alloy; Solid solution; Segregation; Grain boundary; Scanning transmission electron microscopy 1. Introduction The pronounced anisotropy of the mechanical proper- ties in addition to the limited room temperature formability and sensitivity to corrosion are recognized as main factors preventing wider use of magnesium alloys in various applications [1]. The sharp crystallographic texture that develops during forming operations influences the physical properties and determines anisotropic behavior of Mg alloys. A significant improvement in properties can be achieved through texture control. New strategies to achieve such control have focused on using additions of rare-earth (RE) elements as strengthening constituents and texture stabilizers. It has been suggested that the addition of RE elements is the most promising approach for weakening the texture in Mg alloys [2–6]. The influence of RE elements on texture development is attributed to various effects, the main one being the recrys- tallization process, which is strongly affected by the pres- ence of RE elements. It has been suggested that the particle-stimulated nucleation governs the recrystallization and contributes to the development of a weaker and more random texture in alloys with a high content of alloying elements [7–9]. Alternatively, for dilute alloys, the results suggest that a weaker texture is a consequence of preferen- tial grain nucleation in the volumes of the material where localized deformation takes place in the form of shear bands [10–12]. In solid solutions, solute-driven effects mod- ify the texture since solute drag is known to influence both the grain boundary (GB) mobility and the recrystallization kinetics [7,13–16]. The large difference in atomic size between RE and Mg atoms results in a high tendency for RE elements to segregate at grain boundaries [15,17]. The atomic size of RE elements also correlates to their solubil- ity limits in Mg, which reflects the susceptibility of solute atoms to produce a drag effect on a boundary [17]. It has http://dx.doi.org/10.1016/j.actamat.2014.06.004 1359-6454/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Corresponding authors. Tel.: +1 9055259140x23498; fax: +1 9055212773. E-mail addresses: [email protected] (M. Niewczas), gbotton@ mcmaster.ca (G.A. Botton). www.elsevier.com/locate/actamat Available online at www.sciencedirect.com ScienceDirect Acta Materialia 79 (2014) 66–73

Segregation and clustering of solutes at grain boundaries in Mg–rare earth solid solutions

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www.elsevier.com/locate/actamat

ScienceDirect

Acta Materialia 79 (2014) 66–73

Segregation and clustering of solutes at grain boundariesin Mg–rare earth solid solutions

M. Bugnet a, A. Kula a,b, M. Niewczas a,⇑, G.A. Botton a,⇑

a Department of Materials Science and Engineering, McMaster University, 1280 Main Street West, Hamilton, Ontario L8S 4L7, Canadab AGH-University of Science and Technology, Cracow, Poland

Received 30 January 2014; received in revised form 27 May 2014; accepted 2 June 2014

Abstract

The present study validates the previously reported investigations about segregation of rare-earth (RE) elements at grain boundariesin Mg–RE alloys and ultimately provides a direct visualization of the distribution of the solute atoms in the structure of a Mg–Gd alloy.It is demonstrated that Gd forms a solid solution within the Mg matrix in addition to substantial segregation at high-angle grain bound-aries in the form of 1–2 nm clusters, with a postulated face-centered cubic Gd structure. The results suggest significant implications forthe texture development during alloy processing and recrystallization, and thus for the mechanical behavior and properties of Mg–REalloys.� 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Magnesium–rare earth alloy; Solid solution; Segregation; Grain boundary; Scanning transmission electron microscopy

1. Introduction

The pronounced anisotropy of the mechanical proper-ties in addition to the limited room temperature formabilityand sensitivity to corrosion are recognized as main factorspreventing wider use of magnesium alloys in variousapplications [1]. The sharp crystallographic texture thatdevelops during forming operations influences the physicalproperties and determines anisotropic behavior of Mgalloys. A significant improvement in properties can beachieved through texture control. New strategies to achievesuch control have focused on using additions of rare-earth(RE) elements as strengthening constituents and texturestabilizers. It has been suggested that the addition of REelements is the most promising approach for weakeningthe texture in Mg alloys [2–6].

http://dx.doi.org/10.1016/j.actamat.2014.06.004

1359-6454/� 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights r

⇑ Corresponding authors. Tel.: +1 9055259140x23498; fax: +19055212773.

E-mail addresses: [email protected] (M. Niewczas), [email protected] (G.A. Botton).

The influence of RE elements on texture development isattributed to various effects, the main one being the recrys-tallization process, which is strongly affected by the pres-ence of RE elements. It has been suggested that theparticle-stimulated nucleation governs the recrystallizationand contributes to the development of a weaker and morerandom texture in alloys with a high content of alloyingelements [7–9]. Alternatively, for dilute alloys, the resultssuggest that a weaker texture is a consequence of preferen-tial grain nucleation in the volumes of the material wherelocalized deformation takes place in the form of shearbands [10–12]. In solid solutions, solute-driven effects mod-ify the texture since solute drag is known to influence boththe grain boundary (GB) mobility and the recrystallizationkinetics [7,13–16]. The large difference in atomic sizebetween RE and Mg atoms results in a high tendency forRE elements to segregate at grain boundaries [15,17]. Theatomic size of RE elements also correlates to their solubil-ity limits in Mg, which reflects the susceptibility of soluteatoms to produce a drag effect on a boundary [17]. It has

eserved.

M. Bugnet et al. / Acta Materialia 79 (2014) 66–73 67

been reported that RE elements with relatively high solu-bility in bulk Mg have a higher tendency to segregate atgrain boundaries [15]. RE elements are known as “slowlydiffusing elements” and are expected to significantly sup-press the GB mobility. Slow grain growth during therecrystallization process reported by some authors maybe explained by this phenomenon, which gives additionalsupport to the solute-drag hypothesis and to the influenceof RE elements on recrystallization of Mg–RE alloys[10,13].

Although the segregation of RE dopants at grainboundaries has been suggested in the literature, and morerecently has been proven by energy dispersive X-ray spec-troscopy (EDS) measurements in a transmission electronmicroscope on Mg–Y and Mg–Gd alloys [14,16,18], thedistribution of RE atoms in the matrix at the grain bound-aries is not known. A direct visualization of dopants in thematerial structure is imperative for understanding the influ-ence of RE elements in Mg–RE alloys.

In the studies of light metals doped with heavier RE ele-ments, high-angle annular dark-field (HAADF) imaging,or “Z-contrast” technique in the scanning transmissionelectron microscope (STEM), is of particular interest.The recent developments of aberration correctors andhigh-brightness electron guns provide access to Z-contrastimaging at the sub-Angstrom level [19–21]. Several studieshave reported the use of high-resolution HAADF–STEMfor investigating the atomic structure of precipitatesformed in Mg–RE alloys after aging, such as Mg–Nd[22,23], Mg–Dy [24], Mg–Y [25], Mg–Sm [26] and Mg–Gd alloys [27,28]. This chemical imaging technique pro-vides a powerful tool to reveal the segregation of heavy ele-ments at grain boundaries [29–31]. Recently, Nie et al. havedemonstrated the ordered segregation of solute atoms intwin boundaries in an Mg–Gd alloy after compression tests[31].

The objective of this paper is to provide quantitativeresults relevant to the distribution of RE elements in thestructure of Mg–RE alloys. Aberration-correctedHAADF–STEM is used to characterize a high-purityMg–Gd solid solution at the atomic level. This approachallows the visualization of the distribution of atoms in solidsolutions within the Mg matrix, provides direct evidence ofsolute segregation at the grain boundaries, and highlightsthe detailed structure of segregating species at GBs.

2. Experimental

A Mg–0.28 at.% Gd binary alloy was produced in labo-ratory conditions, under argon atmosphere, by meltingappropriate amounts of high-purity Mg and Gd elementsin a spectroscopically pure graphite crucible in an induc-tion furnace, and further casting into a stainless steel moldand cooling in iced water.

The resulting ingot was homogenized for 72 h at 550 �Cunder argon atmosphere. After homogenization, the ulti-mate composition of the alloy was determined by

inductively coupled plasma mass spectrometry; no impuri-ties were detected in the alloy.

The ingot was rolled to a thickness of �3 mm at roomtemperature. The rolling schedule involved multiple rollingpasses combined with annealing to produce recrystallizedmaterial for the next rolling pass. The thickness reductionper pass was controlled not to exceed 20% during the coldrolling.

The samples were machined from a rolled sheet andannealed at 400 �C for 30 min to produce the recrystallizedmicrostructure. A high-purity (99.995%) Mg single crystal,produced by a modified Bridgeman technique, was studiedfor comparison purposes.

A Hitachi SU-70 scanning electron microscope (SEM)was used to analyze the microstructures of annealed mate-rials under channeling contrast conditions. Samples forSEM and electron channeling contrast imaging (ECCI)observations were polished using SiC paper up to 4000mesh, followed by mechanical polishing with a 6, 3, 1 lmwater-free diamond suspensions and final polishing usinga 0.05 lm colloidal silica suspension. ECCI was conductedat acceleration voltages between 5 kV and 20 kV, using asolid-state five-quadrant BSE detector. Working distancesof 7–10 mm were used with low-tilting sample configura-tions. The microscope was operating in the “high current”mode and an objective lens aperture of 50 lm was used.The experimental electron beam parameters were the fol-lowing: probe current: 2–12 nA; and beam convergenceangle: 3–7 mrad.

Electron transparent foils suitable for STEM observa-tions were prepared by the standard procedure, with thefinal electropolishing of 130 lm thick disks carried out withStruers TenuPol-5 twin-jet in a solution containing2.5 vol.% perchloric acid in ethanol. After electropolishing,the specimens were cleaned in a series of ethanol baths. Thefoils were held in ultra-high vacuum for 12 h at 100 �C andthen plasma-cleaned for 5 min to avoid contaminationunder the electron beam. It is important to mention thatSTEM observations before and after 100 �C/12 h annealingand subsequent plasma cleaning were performed, and indi-cated that these steps did not induce structural modifica-tions of the GBs and the grain interior. Atomic-resolution STEM studies were performed in an aberra-tion-corrected (probe-forming lens) FEI Titan Cubedmicroscope operating at 300 kV and equipped with aGatan GIF Quantum electron energy-loss spectrometer.No evidence of damage from the electron beam wasobserved during the experiments. The STEM images pre-sented in this work have been acquired using a FischioneHAADF detector with a convergence angle of �19 mradand a collection angle in the range of 64–200 mrad.

3. Results and discussion

The microstructures of the annealed Mg and the Mg–0.28 at.% Gd alloy are shown in Fig. 1a and b, respectively.Equiaxed and recrystallized grains are formed during

Fig. 1. SEM electron channeling observations of the microstructure of (a)pure Mg and (b) Mg–0.28 at.% Gd annealed at 400 �C for 30 min.

68 M. Bugnet et al. / Acta Materialia 79 (2014) 66–73

annealing in both materials. The effect of Gd on the grainsize of the recrystallized alloy is clearly visible in thesemicrographs. The observed grain size of pure Mg is�200 lm, whereas the Mg–0.28 at.% Gd alloy shows agrain size of �40 lm, suggesting a strong suppression ofthe grain growth due to the Gd atoms. The effect of grainrefinement is enhanced by increasing the Gd concentrationin the solution [32].

The HAADF–STEM images in Fig. 2a and b showthe atomic structure of an Mg single crystal and anMg–0.28 at.% Gd alloy, respectively, observed along the

Fig. 2. Solid-solution alloying with Gd in Mg. HAADF–STEM images ofmagnification of the area highlighted by a square in (b); (d) EDX spectrum ofblur using the Gatan Digital Micrograph software.

[00 01] zone axis. The characteristic hexagonal arrange-ment of Mg atoms is clearly resolved, while Gd dopantsare detected as significantly brighter “dots” in some ofthe atomic columns. The intensity of the atomic columnsin pure Mg is quite uniform, while, in the Gd-doped Mg,there is a very significant variance in the intensity of the lat-tice. In some areas of the microstructure, hexagons of theMg matrix that possess 2–3 Gd-rich columns are observed,as shown in Fig. 2b and c, but Gd atoms are mostly dis-persed in the solution as quasi-randomly isolated atoms.Overall, the solid-solution contains no, or very limited,Gd precipitates distributed in the matrix. Additional obser-vations were carried out in several grains of the Mg–0.28 at.% Gd sample, leading to nearly identical resultsand the detection of a very limited amount of Gd-rich pre-cipitates in the Mg matrix.

EDS, performed within the matrix, away from GBs,confirmed the presence of solute atoms as shown inFig. 2d. The detection of fluorine and chlorine is attributedto residues from the solution used during the electropolish-ing procedure of the samples, whereas carbon and oxygenarise from sample contamination under the electron beamand unavoidable surface oxidation of the thin foil duringspecimen handling, respectively.

Grain boundaries influence many of the structural andphysical properties of engineering materials. In the particu-lar case of doped materials, one of the major concerns is thesegregation of solute atoms at GBs. Fig. 3a shows a high-angle GB in the Mg–0.28 at.% Gd sample. The bright inten-sity in the HAADF–STEM signal visible at the GB can beattributed to the significant segregation of Gd. This effecthas been confirmed by electron energy-loss spectroscopy(EELS), as discussed later. Fig. 3b and c show images ofthe same area, with preferential orientation of the two neigh-boring grains adjacent to the boundary, along [0001] for thetop-right grain in Fig. 3b, and h11�23i for the bottom-leftgrain in Fig. 3c. Independently of the orientation of thegrains, the segregation is not homogeneous along the GB,but instead Gd atoms tend to form crystalline clusters of�1–2 nm in size. Gd clusters are clearly resolved when thegrain on the left or right side of the GB is oriented along amajor zone axis. In spite of a tilt of a few degrees in crystal

an (a) undoped Mg single crystal, (b) Mg–0.28 at.% Gd alloy and (c)the Mg–0.28 at.% Gd alloy. Images (a–c) were smoothed with a Gaussian

M. Bugnet et al. / Acta Materialia 79 (2014) 66–73 69

orientation between the images shown in Fig. 3b and c, theobservation of Gd-rich areas along two different orienta-tions suggests the absence of rod-shaped precipitates alongthe GB. Instead, clusters are observed within the GBs. Fur-thermore, Gd atoms are detected in individual atomic col-umns of the Mg matrix away from the GB, in the grainoriented along the [00 01] zone axis (see top-right grain inFig. 3b). Single Gd atoms are also detected on the other sideof the GB, at the surface of the non-oriented grain (see bot-tom-left grain in Fig. 3b and on the top-right grain inFig. 3c). The observation of Gd-rich clusters was performedon several other high-angle GBs, as illustrated with theexample in Fig. 3d, leading to similar conclusions. The dif-fuse bright HAADF–STEM intensity observed close to theGB, on the bottom-left grain in Fig. 3b, arises most likelyfrom the GB being oriented not exactly edge-on along thedirection of the electron beam. Under such conditions theburied, out-of-focus, Gd-rich clusters at GB are expectedto contribute to the asymmetry in the intensity. This short-range effect is also visible in Fig. 3c and d. It is suggested thatthis kind of diffuse intensity comes from the geometry ofHAADF and the Z-contrast signal that is collected. There-fore, the degree of coherency of the GB, while it cannot beexcluded as a source of the diffuse intensity, is expected tohave a minor influence on the image in this case. The long-range differences in intensity on the two sides of the bound-ary, visible in Fig. 3a, can be explained by the differences inorientation of the two grains. The grain on the left side is noton-axis, in contrast to the grain on the right side, and chan-neling on the atomic columns is not as effective.

The elastic strain fields in the specimen can also modifythe contrast in STEM-ADF imaging in a complex manner,depending on the foil relaxation, the detector angle and thethickness of the investigated area [33–37]. Although theelastic strain contrast in STEM-ADF is particularly visiblefor light elements, such as B dopants in Si [33,34], its con-tribution to the image in the presence of a heavy element,such as Gd, is expected to be small compared to the chem-ical contrast itself, as observed in the case of Er in SiC [38].Therefore, the elastic strain contrast due to the Gd-richclusters might contribute to the bright diffuse band

((b) (a)

20 nm 2 nm

Fig. 3. HAADF–STEM images of GBs in an Mg–0.28 at.% Gd alloy. The ima(c) correspond to the same area of a GB viewed in different orientations of thethe specimen compared to (a–c).

observed around the GBs, but its contribution is likely tobe small compared to the Z-contrast from the Gd atoms.

The segregation of Gd atoms in the form of Gd-richclusters at high-angle GBs studied here shows a differentatomic arrangement than the periodic segregation of soluteatoms in a coherent twin boundary observed in a deformedMg–0.2 at.% Gd alloy [31]. However, the present work andthe study in Ref. [31] do not permit a detailed comparisonto be made because the processing and the state of thematerials were different and the coherency of the GB inFig. 3b–d has not been determined.

Recent studies of diffusion couple in Mg–Gd systemshave shown that diffusivity of Gd in Mg is anisotropicand it is higher along the a-axis than the c-axis [39]. Thetemperature-dependent diffusion coefficients of Gd alonga- and c-axes have been given by approximate equationsD = 1.27 � 10�9exp(–79,268/RT) m2 s�1 andD = 1.79 � 10�9 exp(�81,687/RT) m2 s�1, where R is thegas constant and T is the absolute temperature. The meandiffusion distance

ffiffiffiffiffi

Dtp

of Gd atoms during 12 h annealingof the sample at 100 �C is �20 nm and could account forsome clustering of Gd atoms spaced an average of �5 nmin the ideal solution of Mg–0.28 at.% Gd alloy. This pro-cess should be more effective during 30 min annealing ofcold-rolled samples at 400 �C where the diffusion distanceis �1.3 lm, and Gd diffusion can be accelerated by thepresence of dislocations and vacancies accumulated in themicrostructure. Therefore, annealing of the sample at400 �C may affect some larger volume of the material closeto the GBs, where Gd atoms should be able to enrich theGBs from a distance �1–2 lm. The present HAADF workindicates, however, that the concentration of Gd inside theGBs is homogeneous and no concentration gradient hasbeen detected between the areas next to the GBs and graininteriors. This result suggests that the segregation of Gd tothe GBs must have occurred already during 12 h homoge-nization of the as-cast ingot at 550 �C; in which case, thediffusion distance of Gd is �55 lm, long enough to allowtransport of Gd atoms to the GBs.

EELS analyses, in combination with high-resolutionHAADF–STEM, were carried out to determine the nature

c) (d)

5 nm 5 nm

ge in (a) is an overview, whereas (b–d) are high-resolution images. (b) andtwo grains adjacent the boundary; (d) corresponds to a different region of

electron probe

Mg K (a) (b)

(c)

Gd M4,5

Fig. 4. HAADF–STEM observations and EEL spectra at a high-angleGB in Mg–0.28 at.% Gd alloy. (a) Survey ADF image; (b) Summed EELspectra in areas corresponding to the bottom left grain (green dot in (a),bottom curve in (b)), at the GB (red dot in (a), middle curve in (b)), andthe top right grain (blue dot in (a), top curve in (b)); (c) ADF intensityacross the GB simultaneously acquired with the EEL spectra. (Forinterpretation of the references to colour in this figure legend, the reader isreferred to the web version of this article.)

70 M. Bugnet et al. / Acta Materialia 79 (2014) 66–73

of segregating species at GBs and their structure in theMg–0.28 at.% Gd alloy. A line scan was performed acrossthe GB in Fig. 4a, in an energy range covering the Gd M4,5

and Mg K edges. In Fig. 4b, EEL spectra shown in greenand blue, corresponding to the bulk material on either sideof the GB, reveal the effective doping of the Mg–Gd alloy,with the detection of the Gd M5 edge at �1181 eV. Aroundthe GB (red spectrum in Fig. 4b), the Gd M4,5 edge is moreintense and confirms the substantial segregation of Gd atthe GB. The fine structures highlighted by arrows, andobserved at �1192 eV, �1202 eV and �1222 eV, arise frommultiple scattering due to the large projected thickness ofthe specimen in the area of investigation, preventingdetailed interpretation of the valence state of the Gd atthe GB.

The HAADF intensity profile acquired simultaneouslywith the EEL spectra is plotted in Fig. 4c. The backgroundand the intensity variations on the right and left are very dif-ferent; this effect is attributed to the orientation along the[0001] zone axis of the grain on the right side. Furthermore,the HAADF intensity profile is asymmetric, with a tailtowards the left side of the GB. This additional intensitymay arise from different channeling conditions on eitherside of the GB due to the preferential orientation of theright side grain, and, most important, from the fact thatthe GB might not have been oriented perfectly edge-onalong the direction of the electron beam, as discussed above.This effect is also visible as a diffuse bright HAADF–STEMintensity on the left side of the GB in Fig. 4a.

It should be noted that the concentration of Gd ispeaked at the GB, and is therefore much higher than inthe matrix. A precise quantification of the Gd contentwas, however, not possible due to the following: firstly,the electron beam did not travel exclusively along theGB, but rather probed the Mg matrix together with theGB; and secondly, Gd-rich clusters are observed, ratherthan a uniform distribution of solute at the GB. Further-more, variations of solute concentration in different typesof GBs are expected to occur, and further studies arerequired to quantify these effects. Nevertheless, the presentresults suggest that the segregation of Gd is large, and theintensity of the Gd M4,5 edge, as compared to the Mg K

edge, indicates that the Gd content at the GB is muchhigher than the nominal concentration of Gd in the alloy.

EELS investigations were performed at the O K edge(not shown) across multiple GBs and show minor surfaceoxidation of the TEM specimen, but no preferential oxida-tion at GBs. In addition, the near-edge structures of the OK edge are essentially characteristic of MgO. Based onthese measurements, a preferential oxidation of Gd at theboundary is ruled out and a structural model for the Gd-rich clusters can be proposed, as discussed later in thispaper. However, minor oxidation of the clusters, whichcould not be detected here, cannot be excluded, andremains an open question that calls for furtherinvestigations.

The segregation of Gd and the formation of Gd-richclusters are clearly shown by the chemical contrast fromthe HAADF–STEM images. The HAADF–STEM imagesof Gd-rich clusters formed at GBs provide additionalinsight into their crystal structure. For the clusters orientedalong a high-symmetry zone axis, as shown by red circles inFig. 5a and b, the atomic arrangement revealed byHAADF–STEM is consistent with a cubic structure viewedalong the h100i zone axis, with Gd atoms observed in allcolumns. For these reasons, the orthorhombic structuresof the Mg7RE phase proposed for b0 precipitates [27,40],as well as the Mg3RE DO19 hexagonal superlattice [40]and the Mg6RE structure [22] suggested for b” precipitates,can be excluded. However, the Mg–Gd binary phase dia-gram [41] shows several other stable compounds, includingMg5Gd [42], Mg3Gd [43], Mg2Gd [44] and MgGd [45]; allof these phases have a cubic crystal structure (see Fig. 5d).It should be mentioned that Das et al. [39] reportedrecently the discovery of a new intermetallic compoundwith Mg4Gd6 stoichiometry, obtained during diffusion cou-ple experiments on the Mg–Gd system. This phase couldnot be considered in the analysis of the HAADF–STEMimages presented here due to the unavailability of struc-tural parameters for this compound. According to theMg–Gd phase diagram [41,44], bulk Gd crystallizes in ahexagonal structure, but exists in a body-centered cubic(bcc) structure above 700 �C. Furthermore, it has beendemonstrated that Gd nanoparticles can be stabilized witha face-centered cubic (fcc) crystal structure when they arebelow a critical size, typically 5–10 nm [46,47], which is

M. Bugnet et al. / Acta Materialia 79 (2014) 66–73 71

larger than the dimensions of the GB clusters observed here(1–2 nm).

As mentioned earlier, a precise quantification of the Gdcontent in the clusters cannot be obtained from the currentEELS analysis due to the 3-D nature of the analyzed areaand the size of the clusters, thus preventing the discrimina-tion of various compounds based on their chemical compo-sition. However, atomically resolved HAADF–STEMimages allow to directly compare the observations with ahard spheres model of the crystal structures of thesephases. The crystal structures projected along the h100izone axis of Mg5Gd, Mg3Gd, Mg2Gd, MgGd and fcc Gdare represented in Fig. 5d. Mg5Gd can be excluded becausethe clusters are smaller than the unit cell (a = 22.344 A[42]), and the HAADF–STEM contrast observed in theclusters circled in Fig. 5a and b does not match the patternof the Gd-rich columns in Mg5Gd. Furthermore, the

(d)

(a)

2 nm 2 nm

Mg5GdMg3Gd

Mg2Gd

3.663 Å

3.0

d1

d1

Gd Mg

(b)

Fig. 5. Gd-rich clusters at high-angle GBs and structural models for possHAADF–STEM images presented in Fig. 3b and c, highlighting Gd clusters (cMgGd [45] compounds viewed along h100i, as well as fcc Gd [51] viewed alonatomic arrangements in terms of geometry and inter-atomic spacings are circledmain text for details). The cluster circled in blue in (c) is well reproduced by tvisualized using the VESTA program [52]. (For interpretation of the referencesthis article.)

atomically resolved HAADF–STEM images shown inFig. 5a–c allow us to measure interatomic distances as adirect means to identify the possible phases. In particular,the interatomic distances marked in Fig. 5a were measuredas d1 = 2.77 ± 0.16 A (this value was confirmed by the clus-ter observed in Fig. 5b). A direct comparison with the the-oretical values for the different Mg–Gd phases viewedalong h100i shown in Fig. 5d excludes the presence of Mg3-

Gd and MgGd precipitates at the boundary because theselower symmetry structures would yield a larger apparentinteratomic spacing. The interatomic distance of Gd atomsin Mg2Gd is only slightly larger and should not beexcluded, even though it does not fit within the experimen-tal error. When considering the accuracy of the here-reported measurements, the experimental Gd–Gd distancesare consistent with the fcc Gd structure of the clusters,suggesting that pure fcc Gd is more likely than Mg2Gd

(c)

1 nm

<100>

MgGd

32 Å

3.818 Å2.755 Å

<110>

<100>

fcc Gd

3.896 Å

3.374 Å

d3

d2

ible GB precipitate phases. (a–c) Magnification of the high-resolutionircles); (d) structural models of Mg5Gd [42], Mg3Gd [43], Mg2Gd [44], andg h100i and h110i. The unit cells are shown as straight lines. Comparable

in red in (a) and (b), suggesting a cubic structure viewed along h100i (seehe fcc structure in the h110i orientation. The structural models in (d) areto colour in this figure legend, the reader is referred to the web version of

72 M. Bugnet et al. / Acta Materialia 79 (2014) 66–73

to stabilize in the GB. In addition, the atomic arrangementand interatomic distances (d2 = 3.18 ± 0.23 A, d3 = 3.85 ±0.23 A) displayed in Fig. 5c are in good agreement with thefcc Gd structure observed along h110i, as shown inFig. 5d. The bcc crystal structure of Gd viewed alongh100i also matches the clusters observed in Fig. 5a and bwith the inter-column distance d1 [48], but it is not consis-tent with the atomic structure of the cluster in Fig. 5c.Finally, it should be noted that the experimental structuralpatterns in Fig. 5a–c could not be matched to the structureof the binary compounds along other major zone axes suchas h110i and/or h111i. In summary, the findings presentedhere are consistent with the stabilization of pure Gdclusters with a cubic lattice, presumably fcc, at GBs inthe Mg–0.28 at.% Gd alloy.

Thermodynamic calculations of the Mg–Gd phasediagram [39,49] show that the enthalpy of formation ofMg–Gd intermetallic compounds decreases in the orderof increasing Gd amount in the compound such thatMg5Gd has higher enthalpy of formation than Mg3Gdfollowed by Mg2Gd and MgGd. The reaction is moreexothermic for higher Gd content compounds and oneexpects that these compounds should be formed first. Thephase diagram indicates that MgGd is formed at theconcentration of Gd of �50 at.%, whereas Mg5Gd isformed at �16 at.% Gd [39,49]. HAADF and EELS dataindicate that the local concentration of Gd in the GB isabove 50 at.% Gd, sufficient to form Mg–Gd compoundsmentioned above; however, the analysis of the clusterssuggests that Gd agglomerates in a pure form. This effectmight be related to the competition between interfacialand elastic strain energy, and the effect of these energyterms on the critical size for the nucleation of precipitates,but more studies are required to better understand theseprocesses.

The results presented here support the hypothesis of thestrong influence of RE elements on the texture develop-ment in Mg–RE alloys. It has been suggested that the seg-regation of RE elements at GBs influences recrystallizationkinetics by retardation [16]. As a consequence, a weakeningof the texture by changing the GB mobility and the promo-tion of nucleation of new grains with a wide range of ori-entations at shear bands and GBs have been proposed[7,10,16]. Texture measurements performed on annealedMg–Gd alloys confirm the texture weakening phenomenonin comparison with the texture of pure Mg processed thesame way as Mg–RE alloys [32]. In addition, it has beenfound that the RE elements studied here are extremelyeffective in reducing tension–compression asymmetry, asdiscussed in further detail in Ref. [32]. It is generally agreedthat tension–compression asymmetry is attributed to theenhanced activity of the extension twin system under com-pression, but not under tensile deformation (e.g. Refs.[8,9]). The recrystallization texture characterized by weakerintensity of basal component and the broader distributionof the basal poles results in more balanced twinning andslip activity during compression, which in turn reduces

tension–compression asymmetry. Another factor affectingtension–compression asymmetry is the grain size refine-ment by the addition of RE elements [50]. In the presentstudy, an effective reduction of the grain size in an Mg–Gd alloy is observed, indicating the strong effect of Gdon grain refinement. This effect can be understood in termsof the solute drag that Gd atoms exert on the moving GBs.The present results reveal that the Gd segregating speciesare stabilized in the form of Gd-rich clusters and are com-monly observed at various high-angle GBs in Mg–Gd. Sim-ilar observations were performed at high-angle GBs in anMg–0.36 at.% Sm alloy (not shown), revealing the segrega-tion of Sm in substantial quantities under the form of Sm-rich clusters of �1–2 nm in size. The addition of Sm to pureMg also results in the reduction of grain size, although thesize reduction is not as significant as in the case of Gd addi-tion, the average grain size determined for the Mg–0.36 at.% Sm alloy being �80 lm. These clusters shouldinfluence the texture development and the recrystallizationprocess. However, the fundamental mechanisms by whichthese clusters may interact with the recrystallization frontsmight be very different than in the case of homogeneouslydistributed solute atoms within the GBs. These aspects ofthe present results must be investigated further. Finally, itis worth mentioning that refinement of the grain size resultsin enhanced strength and ductility of Mg–RE alloys, whichis the other benefit of using RE as an alloying element inMg-based alloys.

4. Conclusions

The distribution of solute elements in the structure of anMg–0.28 at.% Gd binary alloy has been studied byHAADF–STEM and EELS techniques. The results revealsignificant segregation of RE solute atoms at high-angleGBs in addition to their quasi-random distribution in theMg matrix. At the GBs, Gd atoms form clusters of �1–2 nm size, with a postulated fcc Gd structure. The resultssuggest that the segregation of solute atoms at GBs canbe directly correlated to the decrease of the grain size inMg–Gd alloy, as compared to pure Mg. These experimen-tal findings provide new insights into understanding theeffect of RE elements on the mechanical behavior and therecrystallization process of Mg–RE solid solutions.

Acknowledgements

The authors are grateful to Steffi Woo for usefulcomments on the manuscript. SEM observations wereperformed at AGH-University of Science and Technologyin Cracow, Poland. The STEM and EELS work wascarried out at the Canadian Centre for ElectronMicroscopy, a national facility supported by the NaturalSciences and Engineering Research Council of Canada(NSERC) and McMaster University, Hamilton, Canada.GAB and MN are grateful to NSERC for financialsupport.

M. Bugnet et al. / Acta Materialia 79 (2014) 66–73 73

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