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Micron 69 (2015) 35–42

Contents lists available at ScienceDirect

Micron

j ourna l ho me page: www.elsev ier .com/ locate /micron

dentifying suboxide grains at the metal–oxide interface of a corrodedr–1.0%Nb alloy using (S)TEM, transmission-EBSD and EELS

ing Hua,∗, Alistair Garnerb, Na Nic, Ali Gholiniab, Rebecca J. Nichollsa,ergio Lozano-Pereza, Philipp Frankelb, Michael Preussb, Chris R.M. Grovenora

Department of Materials, University of Oxford, Parks Road, Oxford, UKMaterials Performance Centre, School of Materials, University of Manchester, Manchester, UKDepartment of Materials, Imperial College London, Royal School of Mines, London, UK

r t i c l e i n f o

rticle history:eceived 3 September 2014eceived in revised form 20 October 2014ccepted 20 October 2014vailable online 29 October 2014

a b s t r a c t

Here we report a methodology combining TEM, STEM, Transmission-EBSD and EELS to analyse the struc-tural and chemical properties of the metal–oxide interface of corroded Zr alloys in unprecedented detail.TEM, STEM and diffraction results revealed the complexity of the distribution of suboxide grains at themetal–oxide interface. EELS provided accurate quantitative analysis of the oxygen concentration acrossthe interface, identifying the existence of local regions of stoichiometric ZrO and Zr3O2 with varying

eywords:irconium alloysS)TEMransmission-EBSDransmission Kikuchi DiffractionELSuboxide

thickness. Transmission-EBSD confirmed that the suboxide grains can be indexed with the hexagonalZrO structure predicted with ab initio by Nicholls et al. (2014). The t-EBSD analysis has also allowedfor the mapping of a relatively large region of the metal–oxide interface, revealing the location and sizedistribution of the suboxide grains.

© 2014 Elsevier Ltd. All rights reserved.

xidation mechanism

. Introduction

Zirconium alloys are used extensively as cladding materials inodern light water reactors. The purpose of the cladding mate-

ial is to separate the uranium dioxide (UO2) fuel rods and theoolant water in order to prevent the escape of fission productshilst maintaining heat transfer to the coolant. The widespreadse of these alloys is mainly due to their low thermal neutron cap-ure cross section and good corrosion performance in the aggressiveeactor environment (Pickman, 1994; Cox, 1961; Cox et al., 1998).

ith increasing demand for high burn-up in modern nucleareactors, environmental degradation of these alloys is now the life-imiting factor for fuel assemblies (Pickman, 1994).

The general waterside corrosion kinetics of zirconium alloysas two stages: the first is an initial pre-transition period of slow

arabolic or cubic corrosion which can be fitted using the power lawquation, �W = Ktn (W is the weight gain in mg/dm2, K dependsn both alloying elements and reactor temperature, and t is thexidation time in days). At a critical thickness, the previously

∗ Corresponding author. Tel.: +44 07583260056.E-mail address: jing.hu@materials.ox.ac.uk (J. Hu).

ttp://dx.doi.org/10.1016/j.micron.2014.10.004968-4328/© 2014 Elsevier Ltd. All rights reserved.

protective oxide begins to break down and there is an abruptincrease in corrosion rate, often termed the transition. The pro-tective oxide then builds up again and post-transition corrosionis composed of several cycles that mimic pre-transition corrosionuntil a period of rapid linear growth eventually develops (Hillner,1977; Garzarolli et al., 2012). Post-transition corrosion generallyfollows an expression of form �W = Kt + C (where C is the initialweight gain). The Zr–1.0%Nb alloy used in this investigation hasbeen shown to exhibit a delayed transition when compared to otherNb-containing alloys (Wei et al., 2012).

During corrosion, an adherent oxide film is formed by the dis-sociation of water molecules at the outer surface of the oxide.Oxygen ions then diffuse inwards, forming an oxide that grows intothe metal. The metal–oxide interface is thus of particular inter-est as it is the location of the oxidation reaction and is expectedto play an important role in controlling the corrosion behaviour.Yilmazbayhan et al. (2006) used Transmission Electron Microscopy(TEM) to show there are 100–150 nm wide rectangular blocky

suboxide grains near the metal–oxide interface on a Zr–2.5%Nballoy. Abolhassani et al. (2010) also identified a substochiomet-ric phase with 40–60 at.% oxygen at the interface of irradiatedZr–2.5%Nb using Energy Dispersive X-ray Spectroscopy (EDS). Niand Hudson (Ni et al., 2012) used Electron Energy Loss Spectroscopy

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EELS) and Atom Probe Tomography (APT) to study oxidisedIRLOTM and Zircaloy-4 alloys, and found a suboxide layer withomposition of ZrO at the metal–oxide interface in pre-transitionamples which disappears after transition. Nishino et al. (1996)sed Auger Electron Spectroscopy (AES) and X-ray Photoelectronpectroscopy (XPS) to study the initial oxidation of Zircaloy-2 atoom temperature, and identified Zr2O, ZrO and Zr2O3 suboxidehases on the oxidised surfaces. The equilibrium ZrO binary phaseiagram contains no stable phases with these stoichiometries, butensity Functional Theory (DFT) modelling performed by Puchaland Van der Ven (2013) and Nicholls et al. (2014) predicts a stablerO suboxide phase with a hexagonal structure as well as other stoi-hiometric phases not in the ZrO binary phase diagram, includingr2O.

It has previously been demonstrated that the micro-chemistrynd crystallography at this metal–oxide interface is extremely com-lex and localised (Ni et al., 2012). The techniques used to examinehe metal–oxide interface are usually focused on a very small regionnd so it is difficult to correlate the local appearance of suboxidehases to the overall corrosion performance of different alloys.

n this paper we present a methodology combining TEM, STEM,-EBSD and EELS analysis to characterise the chemistry and crystal-ography of suboxide phases at the metal–oxide interface. T-EBSDlso allows for characterisation of a relatively large suboxide region,iving a more general view of what is happening at this particulartage of the oxidation process.

. Materials and methods

The sample used for this investigation was prepared from WestinghouseTM developmental alloy with composition ofr–0.9Nb–0.01Sn–0.08Fe (wt%) in the recrystallized condition. Theample was oxidised in an autoclave at EDF Energy under simu-ated Pressurised Water Reactor (PWR) water conditions at 360 ◦Cor 360 days, and shows no sign of accelerated oxidation rate ando is considered to be in the pre-transition stage of the corrosionycle (Wei et al., 2012).

A TEM sample was prepared by in situ lift out method onn FEI FIB 200 using milling currents of 7000–100 pA at 30 kV,nd further thinned down to a thickness below 100 nm in aeiss Nvision 50 dual beam FIB using a beam current of 150 pAt 30 kV and then 250 pA at 5 kV to create a homogeneous andlectron-transparent sample for TEM, STEM and t-EBSD analysis.he sample was further thinned down to ∼50 nm using the low-oltage conditions for final EELS analysis on selected regions ofnterest.

TEM and STEM analysis were performed on a JEOL 2100 LaB6icroscope operated at 200 kV. EELS analysis was performed on

n FEI Titan microscope operated at 300 kV and equipped with aatan image filter. The convergence half-angle was 10 mrad andollection half-angle was 12 mrad with an energy dispersion of.5 eV per channel and a step size of 10 nm. T-EBSD (Trimby, 2012)as performed on an FEI Magellan FEG-SEM XHR 400L at 30 kVith a probe current of 1.6 nA. The step size for t-EBSD was 15 nm,ith an acquisition speed of 33.5 Hz, so that the time to acquire aap time was ∼70 mins. In order to achieve optimal spatial res-

lution, the EBSD measurements were performed in transmissioneometry. The sample was tilted at 20◦ away from the EBSD detec-or with a working distance of 2 mm. The Kikuchi patterns were

ndexed using the AZtec software suite developed by Oxford Instru-

ents. The EBSD measurements in transmission geometry allowedn optimal spatial resolution of ∼10 nm, necessary for the studyf the extremely fine microstructure at the metal–oxide interfaceGarner et al., 2014).

(2015) 35–42

3. Experimental results

3.1. (Scanning) transmission electron microscopy

Fig. 1(a) shows a bright field (BF) image of the TEM sample.Below the Pt layer, which is used to protect the sample during FIBmilling, a ZrO2 layer of thickness 2.8 �m can clearly be seen. Theoxide layer consists of equiaxed grains on the outer surface andcolumnar grains towards the metal–oxide interface. There are somelateral cracks visible throughout the oxide, the largest of whichare located near the metal–oxide interface. The high angle annu-lar dark-field (HAADF) STEM image of the same area in Fig. 1(b)shows metallic second phase particles embedded in the oxidewith voids above them. Below the metal–oxide interface, largezirconium metal grains are visible in both images, with embed-ded metallic second phase particles. These second phase particlesare �-Nb and Zr(Nb,Fe)2 particles. A similar general microstruc-ture has been seen in oxidised Zr alloys by many previous authors(Yilmazbayhan et al., 2006; Ni et al., 2012). The insets in Fig. 1(a)and (b) are higher magnification TEM bright field and STEM HAADFimages of the region chosen for EELS analysis. In all these images,there is a slight contrast change at the metal–oxide interface butit is difficult to distinguish if there are phases other than ZrO2 andmetallic Zr in this region.

Convergent beam electron diffraction (CBED) patterns wereacquired from a region close to the metal–oxide interface justbetween the two cracks shown in the higher magnification TEMbright field image in Fig. 2(a). A suboxide grain which is stronglydiffracting is highlighted by the dotted lines. A STEM dark fieldimage of the suboxide grain area is shown in Fig. 2(b); the Z-contrasthelps to identify the suboxide as a region of intermediate contrastbetween the brighter metal substrate and the darker stoichiomet-ric ZrO2, highlighted by arrows in the figure. The area where theCBED pattern was taken is also highlighted. A typical diffractionpattern is shown in Fig. 2(c) and can be indexed as matching withthe [324] zone axis of the hexagonal ZrO phase with P-62m sym-metry and lattice parameters a = 5.31 A and c = 3.20 A predicted byNicholls et al. (2014).

3.2. Transmission- EBSD

The band contrast image from the t-EBSD analysis shown inFig. 3(a) highlights the typical oxide microstructure, with smallequiaxed grains towards the outer surface and large, elongatedmonoclinic grains towards the interface, oriented parallel with theoxide growth direction (Yilmazbayhan et al., 2006; Ni et al., 2012).At the interface, blocky grains are clearly visible, which correspondwell with the appearance of suboxide grains described by otherauthors (Yilmazbayhan et al., 2006). Some of the grains at theinterface are identified as the hexagonal ZrO phase with P-62msymmetry and lattice parameters a = 5.31 A and c = 3.20 Å (Nichollset al., 2014), while other grains match Zr metal with P63/mmcsymmetry and lattice parameters a = 3.23 A and c = 5.14 Å (Hull,1921). The patterns were also checked against the orthorhombicZrO phase predicted by Nicholls et al. (2014), however there wasno match in any of the interfacial grains. To show how we ver-ify the structure of these suboxide grains, the Kikuchi patterns(EBSPs) from this region were matched using the AZtec softwaresuite with the hexagonal ZrO phase predicted by Nicholls et al.(2014). Examples of matched EBSPs acquired from the suboxideregion are shown in Fig. 4, with the software providing consis-

tent matching with a mean angular deviation of generally ∼0.5◦.The phase map for the entire region is shown in Fig. 3b. Thismap shows, as expected, that the majority of the oxide is com-posed of columnar monoclinic ZrO2 grains, and small tetragonalZrO2 grains are also visible throughout the oxide. The tetragonal

J. Hu et al. / Micron 69 (2015) 35–42 37

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ig. 1. (a) TEM bright field image of the TEM sample. (b): HAADF-STEM image of

agnification images of this same area after further FIB thinning (which has create

hase fraction is ∼1%, and the average grain size of the tetragonalrO2 grains is about 20 nm. This tetragonal phase fraction is lowerhan the value of ∼3% measured on this alloy at a similar expo-ure time using synchrotron X-ray Diffraction (XRD) (Wei et al.,013). This is thought to be due to the transformation of stress-tabilized tetragonal grains to the equilibrium monoclinic phaseuring preparation of the thin FIB sample, also the possibility thathe majority of the remaining small tetragonal grains are belowhe resolution limit of t-EBSD. At the metal–oxide interface, theres a clear layer composed of hexagonal ZrO grains. The ZrO grainsre generally equiaxed and vary in size from ∼50 to 300 nm. Therientation map in Fig. 3(c) indicates that most (but not all) of theuboxide grains inherit the orientation of the metal grain on whichhey form. Several examples of this epitaxial growth of suboxidere circled.

.3. EELS analysis

Several EELS line profiles were performed in the region high-

ighted in Figs. 1–3, and the exact positions of the profiles are shownn the detail of the t-EBSD image in Fig. 5(a) and the HAADF imagen Fig. 5(b). The sample thickness after final thinning for EELS wasstimated to be around 50 nm. The oxygen concentration was cal-brated by using the starting point of the line in the ZrO2 region

M sample. The area used for EELS analysis is highlighted. The inset shows higherhole under the left hand crack).

using the method described by Ni et al. (2011). In this region, both t-EBSD and CBED have confirmed the existence of hexagonal ZrO. TheHAADF image also shows grain structures with different contrastsat the interface, and several second phase particles.

Four EELS profiles were taken in total from this region, with allfour detecting some regions with ZrO stoichiometry of thicknessesvarying from 13 nm to 326 nm (Fig. 6). The difference in thickness ofthe ZrO suboxide regions, even over such a small area, demonstratesthe extremely heterogeneous nature of the suboxide, as has beencommented on in previous work (Ni et al., 2011). In addition, inthe third profile a layer with a Zr:O ratio of 3:2 is detected. Thismay be a result of an inclined interface running through the samplethickness, therefore there is a contribution from overlapping Zr andZrO grains. It has previously been shown that the t-EBSD patternswill be created preferentially from the exiting surface of the TEMsample (Suzuki, 2013), and so it is likely that in this area there ismetallic Zr on the front side of the sample and ZrO on the backsurface.

We can be confident of the accuracy of these compositional

measurements because we measure the ZrO2 stoichiometry on theoxide side of all the scans. On the metal side of the interface thereis a gradient of oxygen into the metal which also seems to varyover the small area analysed here, also in agreement with previ-ous observations (Ni et al., 2012). On line profiles 2 and 3, there

38 J. Hu et al. / Micron 69 (2015) 35–42

Fig. 2. (a) High magnification TEM bright field image of the area near the metal–oxide interface, which was chosen for EELS analysis. The suboxide grain, which was in as age o( al ZrO

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trongly diffracting condition, is highlighted using dotted lines. (b) STEM dark field imc) CBED pattern of the suboxide grain which matched with the predicted hexagon

icholls et al. (2014)

eems to be another plateau region of around 100 nm thicknessfter the suboxide region with an oxygen content of 25 at.% (Zr3O),hich suggests that the Zr metal here is saturated with oxygen. In

ll four line scans, the oxygen content drops to a value close to zeroeeper into the metal, indicating that the final sample cleaning hasuccessfully removed most of the surface oxide.

. Discussion

TEM images, as shown in Figs. 1(a) and 2(a), can be used to high-ight the microstructure of the oxide and the metal-oxide interface,or instance the columnar ZrO2 grain structure and the metal canlearly be distinguished. However, diffraction contrast imaging ofhe suboxide grains does not help us understand their chemistry.TEM HAADF images like those in Figs. 1(b) and 5(b) are more use-ul due to the strong atomic number contrast (Pennycook et al.,996), so that the weight difference from different oxygen concen-rations can easily be visualised. For instance, a suboxide region of

ifferent contrast at the metal–oxide interface can readily be seen

n Fig. 5(b).The t-EBSD maps of the entire FIB sample are shown in Fig. 3.

ue to overlapping patterns at oxide grain boundaries, the oxideicrostructure is highlighted in a way that is not possible with

f the suboxide grain area, with the area where CBED pattern was taken highlighted. phase with P-62m symmetry and lattice parameters a = 5.31 A and c = 3.20 Å.

conventional TEM imaging. In the band contrast image, blockygrains are also visible at the metal oxide interface, and these grainsare indexed with high confidence as the hexagonal ZrO phase pre-dicted by DFT modelling (Nicholls et al., 2014), as shown in thephase map in Fig. 3(b). The matched Kikuchi patterns in Fig. 4confirm that this modelled phase provides a good fit for this struc-ture at the interface. In the region chosen for EELS analysis, thereis a ZrO grain of diameter ∼300 nm, as well as smaller suboxidegrains on either side. This map also shows the variation of thicknessof the suboxide region along the metal–oxide interface. Althoughit appears that the ZrO layer is not continuous from the t-EBSDmap, this could be due to the limited resolution of the technique(∼10 nm) since all of the EELS line profiles show at least a very thinZrO layer at the interface. More EELS profiles would be required toconfirm this.

Because we can precisely identify the positions of the EELS lineprofiles on the HADDF and t-EBSD images, we can discuss howresults from the different the techniques compare on this small

region of the interface. EELS profile 1 shows a relatively small ZrOregion, which agrees well with the t-EBSD map at this location. Thesecond EELS profile is taken across the centre of the large (∼300 nm)diameter suboxide grain identified by t-EBSD as hexagonal ZrO andshows a clear 1:1 ZrO stoichiometry across the entire grain. The

J. Hu et al. / Micron 69 (2015) 35–42 39

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ig. 3. (a) Band contrast map, (b) phase map and (c) orientation map from t-EBSD aBED analysis is also highlighted in (b). Circles in (c) show examples of epitaxial gro

hird profile, which is apparently taken across the edge of the samerain, appears to show a slightly different stoichiometry with a Zr:Oatio of 3:2. It shows a small ∼13 nm region of ZrO, which may coin-ide with the small oxide grain identified above the right hand sidef the large grain in the t-EBSD image. As previously mentioned, it

s possible that this increase in Zr ratio is caused by signal from anverlapping Zr grain, due to the reduced ZrO thickness at the edgef the suboxide grain. The fourth EELS profile again shows the ZrOtoichiometry over a distance of ∼300 nm at the interface. Com-arison with the t-EBSD phase map suggests that the profile goes

s of the TEM sample (grain boundaries >10◦ are highlighted). The area for EELS andf the ZrO suboxide on the underlying Zr grains.

across two suboxide grains with the same hexagonal ZrO crystalstructure. It is also important to note that the suboxide structurevisible in the HAADF image in Fig. 5(b) correlates well with thet-EBSD of this region in Fig. 5(a).

Another interesting outcome of this work is the variation in the

oxygen concentration gradient into the metal after the interfaceregion. A comparison of EELS profiles 2 and 4 (Fig. 6.) shows thatin two regions that show a similar thickness of ZrO, the oxygenconcentration gradients into the metal are quite different. The oxy-gen concentration along line profile 2 reduces to below 10% over a

40 J. Hu et al. / Micron 69 (2015) 35–42

Fig. 4. Selected electron backscattered patterns (EBSPs) from the suboxide region of the TEM sample and corresponding solutions using AZtec software and ZrO hexagonalphase (a) mean angular deviation (MAD) = 0.45◦ number of matched bands = 7, (b) MAD = 0.59◦ number of bands = 7, (c) MAD = 0.57◦ number of bands = 8.

Fig. 5. Positions of EELS line scans from the suboxide region relative to (a) the t-EBSD map. (b) The HAADF image. All profiles started from the oxide towards metal.

J. Hu et al. / Micron 69 (2015) 35–42 41

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ig. 6. Positions of EELS line scans from the suboxide region relative to the HADDF

17 nm ZrO.

istance of ∼100 nm, whereas along profile 4 has still not reachedhis level over ∼300 nm. It can be seen from Fig. 3 that beyondhe suboxide grain, profile 2 enters a single ZrO2 grain, whereasrofile 4 enters a region of smaller ZrO2 grains with a high levelf non-indexed points. It is therefore possible that there are over-apping grains in this region causing different contributions to ther:O ratio along profile 4. It has previously been shown that theresence of the suboxide is correlated with corrosion rate and dis-ppears after transition due to the increased local corrosion rate (Nit al., 2012). As mentioned previously, this sample appears to notave undergone the kinetic transition and so is in the pre-transitioneriod of slow oxide growth, adding further evidence to the con-ept that the presence of the suboxide can be correlated with slowxide growth. We note here that the transition could have occurredocally at some regions of the metal–oxide interface, even prior tohe global increase in corrosion rate.

Comparing with data from other Zr alloys, it appears that thisample (Zr–0.9Nb–0.01Sn–0.08Fe) has a thicker suboxide layerhan faster corroding alloys such as ZIRLOTM. EELS data shows thatre-transition ZIRLO samples contain a 50 nm ZrO layer which ishinner than is observed here (Ni et al., 2011). Care must be takenhen comparing these results as these observations clearly show a

ignificant variation in suboxide thickness along even short lengthsf the interface, but these results suggest that good oxidation resis-ance seems to be related to a thicker suboxide layer. At this pointt is unclear whether this contributes to or is a result of the slowerxidation kinetics.

. Conclusions

Here we report a methodology combining TEM, STEM, EELSnalysis and t-EBSD to study suboxide phases at the metal–oxidenterface of corroded Zr alloys. The appearance of the suboxidehase is shown to be localised and varied in distribution and com-osition, which makes it difficult to visualise in TEM imaging and

herefore to acquire diffraction patterns to identify possible phases.-EBSD gives an easier and more accurate way to match the candi-ate phases over a larger area and clearly reveals a suboxide layert the metal–oxide interface. This suboxide is made up of approx-mately equiaxed grains of varying sizes, and it is clear that they

and t-EBSD map: (a) 20 nm ZrO, (b) 326 nm ZrO, (c) 13 nm ZrO + 259 nm Zr3O2, (d)

do not form a uniform layer at the interface. EELS analysis allowsfor quantitative analysis of the oxygen concentration across theinterface and reveals the presence of stoichiometric ZrO.

This work has shown that it is necessary for the suboxide regionto be examined with a number of complementary techniques toensure accurate characterisation of the complex microstructure.We have demonstrated that a number of techniques can be appliedto exactly the same volume of material if the original TEM specimenis of the appropriate quality and robustness. We believe that com-bining the techniques outlined here can be a powerful methodologyused to characterise both the interfacial chemistry and crystal-lography of a wide range of engineering alloys where a complexnano-structure, often containing non-equilibrium phases stabilisedby stress or other local parameters, may be difficult to study usinga single technique.

Acknowledgements

This project is part of the MUZIC-2 collaboration studying oxi-dation and hydrogen pickup in Zr alloys. The authors gratefullyacknowledge support from EDF Energy, Westinghouse and RollsRoyce. Also, Jing Hu is supported by the China Scholarship Council.

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