Corrosion Science 53 (2011) 3623–3635

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Corrosion Science

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Influence of dissolved oxygen on oxide films of Alloy 690TT with differentsurface status in simulated primary water

Zhiming Zhang, Jianqiu Wang ⇑, En-Hou Han, Wei KeState Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China

a r t i c l e i n f o

Article history:Received 2 March 2011Accepted 1 July 2011Available online 14 July 2011

Keywords:A. NickelB. TEMC. OxidationC. Intergranular corrosionC. Effects of strain

0010-938X/$ - see front matter � 2011 Elsevier Ltd. Adoi:10.1016/j.corsci.2011.07.012

⇑ Corresponding author. Tel.: +86 24 23893723; faxE-mail address: wangjianqiu@imr.ac.cn (J. Wang).

a b s t r a c t

Oxidation of Alloy 690 with different surface conditions were performed in oxygenated primary watercontaining B and Li at 325 �C and 15.6 MPa for 2 months. Oxide scales were analyzed using various meth-ods. Results showed triple-layer oxide films were formed: the outmost layer with dispersed big oxideparticles; the intermediated layer with loose sheet-like or needle-like oxides and the inner layer withincompact cellular oxides, which were lack of protectiveness. Oxide morphologies were affected by thesurface status of samples. Dissolved oxygen increased the intergranular attack (IGA) sensitivity of Alloy690TT. Growth mechanism of the oxide film was discussed.

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1. Introduction

Stress corrosion cracking (SCC) has become one of the main fail-ure mechanisms of steam generator (SG) tubing of pressured waterreactors (PWRs) [1]. Several SCC mechanisms of SG tubing andsome other materials in the simulated water solution of nuclearpower plants have been proposed during the past several decades[2–12]. Although the detailed micro-processes of the crack initia-tion and propagation between these mechanisms are different,the properties of the oxide films (nature, microstructure, protectiveaspect, brittleness, porosity, etc.) always play a crucial role for allproposed mechanisms [2–12]. It is well-known that a compactand continuous surface oxide film could protect the materials fromfurther corrosion [6]. Therefore, oxygen is often added to the pri-mary water to stabilize the surface oxide and avoid effluent of irra-diated oxide during interval inspection and operation at lowtemperature [13]. The primary water may not be totally deaeratedduring start-up and sometimes during operation. The oxygen couldalso be introduced in the primary water by the feed water or by theradiation decomposition of the water at the reactor cores wherethe radiation intensity was very high [14]. How does dissolvedoxygen affect corrosion behavior of Alloy 690TT in primary side?

It had been discussed that the corrosion behavior of materials inhigh temperature water were closely related with the samples’surface statuses and the water chemistry environment (such asthe dissolved oxygen or hydrogen). [4,15–28]. McIntype et al.had found that a much thicker oxide film was formed on Inconel

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600 when oxidized in the water with low concentrations of oxygenpresent and increased oxygen concentration in the water pro-moted the increased dissolution of chromium [18]. Kim suggestedthat a thicker oxide was formed on Type 304 stainless steel (ss) un-der 200 ppb O2 or 200 ppb H2O2 conditions than under 150 ppb H2

and a closely packed oxide with a small particle size was formed onType 316ss under 200 ppb oxygen and 20 ppb hydrogen condi-tions, while a less-packed oxide with a large particle size wasformed under 150 ppb hydrogen and 15 ppb oxygen conditions[19,20]. Kumai and Devine found that the oxidation rate of Fe inoxygenated high temperature water decreased with the increaseof oxygen concentration from 8 to 22 ppb and Fe3O4 was trans-formed to a-Fe2O3 when the dissolved oxygen concentration wasraised from 22 to 53 ppb [21]. The oxide films grown on Ni–Cr–Fe Alloys in high-purity water at 288 �C were a function of the alloychromium concentration and corrosion potential, which was con-trolled by the dissolved oxygen concentration [22]. It suggestedthat dissolved oxygen in high temperature water could affect theoxidation behaviors of Ni-based alloys and stainless steels.

The purpose of this paper is to investigate the effect of dissolvedoxygen on corrosion mechanisms of Alloy 690TT with differentsurface statuses in simulated primary water.

2. Experimental

2.1. Materials and sample preparation

Samples used in the present experiment were cut from thermallytreated (TT) Alloy 690 tubing of 19.07 mm diameter by 1.09 mmwall thickness. Its chemical composition is listed in Table 1. The

3624 Z. Zhang et al. / Corrosion Science 53 (2011) 3623–3635

tubes were cut longitudinally into four parts and then were flattenedinto sheets [29]. These flattened sheets were then thermally treatedat 715 �C for 2 h to release the residual stresses. Alloy 690TT sampleswere then cut into 10 mm by 8 mm from these sheets.

Some samples were mechanically ground with waterproofabrasive paper up to 400 grit. Some were mechanically groundup to 1500 grit. Some samples were ground up to 2000 grit andthen mechanically polished using 2.5 lm diamond paste. Half ofmechanically polished samples were then electropolished using30 Vol% nitric acid (HNO3) + methanol (H3COH) at 25 V for about3 s. All samples were then washed with acetone, ethanol anddeioned water. So samples with four kinds of surface states:ground to 400 grit (ground), ground to 1500 grit (ground), mechan-ical polished (MP) and electropolished (EP), were prepared.

2.2. Immersion test

Alloy 690TT samples were immersed in a dynamic 7L Type 316autoclave with a circulating loop. Deionized water was added intothe loop with ion exchange resin to further purify the water. Onethousand and five hundred milligrams per liter of B as boric acid(H3BO3) and 2.3 mg/L Li as lithium hydroxide monohydrate(LiOH�H2O) were added into the deionized water to simulate theprimary water. Ion exchange resin was excluded from the circulat-ing loop after adding B and Li. The measured pH was about 6.2 atroom temperature and it was consistent with the calculated value.The solution was first bubbled with high-purity nitrogen to controlthe dissolved hydrogen (DH) content to be lower than 0.02 cm3

(STP) H2/kg H2O. Then the dissolved oxygen (DO) in the inlet waterwas kept at 1.40 cm3 (STP) O2/kg H2O by purging the high-purityoxygen into the water tank of the loop. After circulating for severaldays when the water chemistry of the solution became stable, theautoclave loop was pressured to about 15.6 MPa and heated to325 �C. The coupons were exposed in the solution for a period of60 days.

2.3. Characterization of the oxide films

Post-test analyses conducted include surface morphologyobservations and microstructural and compositional analyses of

Fig. 1. SEM photographs of the oxide films on Alloy 690TT with different surface conditio(STP) O2/kg H2O at 325 �C: (a) ground to 400 grit; (b) ground to 1500 grit; (c) mechanic

the oxide films grown on Alloy 690TT with different surface statesusing various methods.

2.3.1. Scanning electron microscopy (SEM) observationMorphologies of surface oxide films were investigated using a

field emission gun SEM (FEG-SEM) Phillip XL30. Chemical compo-sitions of some oxides in the oxide films were analyzed by the SEMwith an energy dispersive X-ray spectroscopy (EDX) system.

2.3.2. Grazing incidence X-ray diffraction (GIXRD) analysisThe oxide phases in the oxide films were analyzed by GIXRD

performed at Shanghai Synchrotron Radiation Facility (SSRF) andNational Synchrotron Radiation Laboratory of University of Scienceand technology China (NSRL). The energy of the photons was ad-justed to 10 keV to reduce sample adsorption. The X-ray incidenceangle was controlled from 0.1� to 0.8� so that the X-ray received bythe detector were mainly came from the oxide films and also in or-der to obtain the phase distribution from the outer layer to the in-ner one. The scan range of the X-ray detectors was from 20� to 55�.Counting time was 1–2 s and step size was 0.018–0.05�. At lastconventional X-ray diffraction measurement was performed onsome samples of which oxide films were very thick. Then the rota-tion angle of the sample plate was from 10� to 28� and at the sametime it was from 20� to 56� for the X-ray detectors.

2.3.3. X-ray photoelectrons spectroscopy (XPS) analysisThe chemical composition of the oxide films by XPS were fin-

ished with an ESCALAB250 using Al Ka radiation (hm = 1486.6 eV)at a pass energy of 50.0 eV. The take off angles of photoelectronswas 45�, with respect to the sample surfaces. For XPS characteriza-tion, Ni2p, Cr2p, Fe2p, O1s and C1s core level spectra had been re-corded. The spectrometer was calibrated against the band energy(BE) of the surface carbon contamination C1s at 285.0 eV. In orderto get the composition–depth profiles by successively removingthe oxide surface with argon ion bombardment, a 2 keV argonion sputtering at a target current of 2 lA/cm2 and a pressure of5.5 � 10�8 mbar was used. The argon ions sputtering rate wasabout 0.2 nm/s. To analyze the individual contribution of theNi2p3/2, Cr2p3/2, Fe2p3/2 and O1s core levels, peak decomposi-tion was carried out with a computer software XPSPEAK4.1 usingGaussian–Lorentzian peak shapes and a Shirley background.

ns after 2 months’ immersion in 1500 mg/L B and 2.3 mg/L Li solution with 1.40 cm3

al polished; (d) electropolished.

Table 1Chemical composition of Alloy 690TT provided by EPRI (wt%).

Element Ni Cr Fe Mn Ti S P C N Si Cu Co Al

wt% 59.50 29.02 10.28 0.30 0.33 0.001 0.009 0.018 0.0234 0.31 0.010 0.015 0.16

Table 2Chemical compositions of big polyhedral oxide particles shown in Fig. 1 by EDSanalysis of SEM.

Ni (at.%) Fe (at.%) Cr (at.%) O (at.%)

On the surface ground to 400 grit 21.80 27.86 03.59 46.93On the surface ground to 1500 grit 25.27 31.26 05.12 38.35On the MP surface 25.98 28.82 04.25 40.95On the EP surface 30.80 36.85 03.47 28.88

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2.3.4. Transmission electron microscopy (TEM) analysisThe surface oxides were scraped using a razor blade and diluted

in ethanol. Then the fine oxides were collected on the surface ofethanol using a Cu grid covered with ultrathin carbon films andanalyzed using a 200 kV Tecnai G2 F20 TEM (FEI, Eindhoven, Neth-erlands), equipped with an EDX system. This EDX system couldn’tanalysis the light elements, so the oxygen content in some oxideswas unknown. The cross-sections of the surface oxide films wereprepared by a FEI QUANTA 200 3D focus ion beam system (FIB)and analyzed using a FEI 300 kV Tecnai G2 F30 high resolutionTEM, equipped with an EDX system.

3. Results

3.1. SEM morphologies of oxide films

As shown in Fig. 1, the morphologies of oxide films were verydifferent on different surfaces after 2 months’ immersion in oxy-genated primary water. The ground (Fig. 1a and b) and MP(Fig. 1c) surfaces were mainly covered with loose sheet-like oxidefilms and an amount of scattered big polyhedral oxide particles ontop. In addition, a few needle-like oxides also spread on these sur-faces. In contrast, EP surface was covered with loose needle-likeoxide films and scattered big polyhedral oxide particles on top,as shown in Fig. 1d. It should be noted that the white boundaryin Fig. 1d was a bundle formed by the end of many needle-like oxi-des and did not correspond to the grain boundary of Alloy 690TT.

The chemical composition of the big oxide particles on differentsurfaces shown in Fig. 1 was analyzed by EDS of SEM, as listed inTable 2. It was clear all big particles were rich in Fe and Ni but de-pleted in Cr.

3.2. GIXRD analysis

The GIXRD results are shown in Fig. 2. For the oxide film grownon the surface ground to 400 grit, the diffraction peaks mainly cor-responded to spinel oxides (such as NiFe2O4, PDF 89-4927; Ni-Cr2O4, PDF 77-0008; FeCr2O4, PDF 89-3855; NiCrFeO4, PDF 52-0068), when the incident angles were 0.1� and 0.2�. This was con-sistent with the corrosion behavior of 304 stainless steel, Alloy 625and 600 in the high temperature hydrogenated water [30–33].However, it was difficult to distinguish these spinel oxides, as thediffraction peak positions of these spinel oxides differ by the de-gree as small as about 0.3�. When the incident angles increasedto 0.5� and 0.8�, the intensity of the diffraction peaks of the spineloxides increased a lot. Moreover, the diffraction peaks of nickeloxides were present at these two higher incident angles. For X-rays

incident on a solid from air, a critical angle is known to exist, belowwhich total external reflection occurs. According to Ziemniak [30],the critical angle for typical spinel oxides (i.e., NiFe2O4, NiCr2O4 andFeCr2O4) was 0.25�. The used X-ray wavelength was 1.24 Å and thepenetration depth of X-ray at the critical angle about 13.8–15 nm.An increase in incident angle from 0.2� to 0.3� increased the sam-pling depth from about 4.3–830 nm and the penetration depthwould be about several micrometers at the incident angle of 0.5�and 0.8�. So the nickel oxides should be located in the inner layerof the surface oxide film. The diffraction peaks of the matrix werestill not very clear, even the incident angle increased to 0.8�, indi-cating the oxide film grown on Alloy 690TT in the present oxygen-ated primary water were very thick. Fig. 2b is the magnification ofthe marked rectangle in Fig. 2a, showing the diffraction peaks withmore details. It could be seen that there were three diffractionpeaks from 34� to 35.5� and the 2h diffraction angles were34.47�, 34.60� and 34.88� respectively. The diffraction angles, 2hof the (1 1 1) diffraction peaks of the nickel oxides, spinel oxidesand matrix are 34.53� (PDF 47-1049), 34.611� (NiFe2O4, PDF 89-4927) and 35.185�(PDF 33-0945) respectively. Therefore Fig. 2bpresented the 2h diffraction profiles of nickel oxide, spinel oxideand Alloy 690TT (matrix).

For the oxide film grown on EP surface, the 2h angles of the dif-fraction peaks, as shown in Fig. 2c and d, were identical with thoseon the ground surface, indicating that the oxides on the EP surfacewere also mainly spinel oxides and nickel oxides.

As a result, GIXRD analyses showed that the oxide films grownon Alloy 690TT in the oxygenated primary water were composed oftwo layers: the outer layer with spinel oxides and the inner layerwith nickel oxides, whatever the surface states and the oxidemorphologies.

3.3. XPS analysis

XPS results of surface films are depicted in Figs. 3 and 4. The ar-gon ion sputtering progress was only performed on the top layer ofthe oxide film on the EP Alloy 690TT because the oxide films onground, MP and EP surfaces were about several micrometers inthickness. The depth profiles of Ni, Cr, Fe and O in the oxide filmon EP surface are shown in Fig. 3. The top layer of oxide film wasrich in Ni, but depleted in Cr severely. The content of Fe was muchhigher than that of Cr. However, Table 2 showed that the outmostbig oxide particles of the oxide film were rich in Fe and the contentof Ni was lower than that of Fe. The sputtering and analyzing areasin the present study were about 2 mm � 2 mm and 0.5 mm �0.5 mm respectively, whereas the distribution of big oxide parti-cles was inhomogeneous (Fig. 1d). The big oxide particles seemto be deposited on the needle-like oxides and might be lost bysputtering due to the weak cohesion. So it was guessed that thechemical composition shown in Fig. 3 might mainly come fromthe needle-like oxides and partially from the big oxide particles.

Considering the GIXRD results and the composition–depth pro-files of the surface oxide film on EP surface, the peak decomposi-tion of the XPS results was carried out, as shown in Fig. 4.

In Fig. 4a, the O2p1/2 peak was systematically decomposed intotwo components for 0 s sputtering: the signal at the band energy(BE) of 531 ± 0.6 eV was assigned to OH� in the M(OH)x, and theone at 530 ± 0.6 eV to O2� in the oxides [34–37]. After 10s and3490s sputtering, the decomposition of O2p1/2 peak was similar

Fig. 2. X-ray diffraction profiles of the oxide films formed on Alloy 690TT after 2 months’ immersion in 1500 mg/L B and 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/kg H2O at325 �C: (a) oxide film on the surface ground to 400 grit; (b) the enlargement of the area surrounded in (a); (c) oxide film on the EP surface; (d) the enlargement of the areasurrounded in (c). Three kinds of arrow mean NiO ( ), Spinel oxide ( ), Alloy 690TT ( ) diffraction peaks.

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with that after 0s sputtering. Considering the sputtering rate ofXPS, the thickness of the oxide film was about 698 nm after

3490s sputtering. It was about several micrometers from thesurface when the nickel oxide was present in the oxide film

Fig. 3. Composition profiles in the depth of oxide films formed on EP Alloy 690TTafter 2 months’ exposure in1500 mg/L B + 2.3 mg/L Li solution with 1.40 cm3 (STP)O2/kg H2O at 325 �C by using XPS.

Z. Zhang et al. / Corrosion Science 53 (2011) 3623–3635 3627

(Fig. 2). So the oxides might mainly correspond to the spinel oxidesdetected by GIXRD. The decomposition of O2p1/2 peak indicatedthat the outer layer of the oxide film was composed of hydroxidesand oxides.

Fig. 4b shows the Ni 2p3/2 core level spectra after differentsputtering time: 0s, 10s and 3490s. The Ni 2p3/2 peaks were sys-tematically decomposed into three main components and the asso-ciated satellites. Compared with published data [18,23,36–38], thepeak at 856.55 ± 0.2 eV with a satellite at 862.75 ± 0.3 eV was as-signed to Ni2+ in Ni(OH)2, the peak at 854.9 ± 0.3 eV with a satelliteat 861 ± 0.5 eV to Ni(II) and the peak at 852.7 ± 0.1 eV with a satel-lite at 858.5 ± 0.2 eV to Ni0. The nickel divalent oxides shouldmainly be the spinel oxides containing Ni detected by the GIXRD,such as NiFe2O4 and NiCr2O4. At the outmost of the oxide film,the element Ni was present as the forms of Ni(II) and Ni(OH)2. Afterabout 10s sputtering, the Ni0 peak appeared. After about 3490s

Fig. 4. The detailed XPS spectra of O1s, Ni2p3/2 and Cr2p3/2 in the oxide film grown on E1.40 cm3 (STP) O2/kg H2O at 325 �C: (a) spectra of O1s; (b) spectra of Ni2p3/2; (c) spect

sputtering, the peak decomposition of Ni2p3/2 peak was similarwith that after about 10s sputtering. Nickel oxide was stable inthe present high temperature primary water with 1.40 cm3 (STP)O2/kg H2O with the consideration of the potential-pH diagram ofNi–H2O at 300 �C [39]. However, nickel oxide was less stable thanother oxides [40–42]: Cr2O3 > FeCr2O4 > NiCr2O4 > NiFe2O4 > NiO.So the detected Ni0 in the oxide film might be reduced from theNiO during the argon ionic sputtering [42].

Fig. 4c shows the Cr2p3/2 core level spectra after different sput-tering time: 0s, 10s, and 3490s. The Cr 2p3/2 peaks were systemat-ically decomposed into two main components: the BE signal of577.6 ± 0.3 eV was attributed to Cr3+ in Cr(OH)3 and the one of576.3 ± 0.3 eV to Cr(III), by comparison with the published data[35,38,43–45]. However, it was suggested that the Cr hydroxideshould be CrOOH [18,46–48]. For Cr(III), it should mainly corre-spond to the spinel oxides containing Cr, such as NiCr2O4 and NiCr-FeO4. The peak decomposition of Cr2p3/2 peak was similar afterdifferent sputtering time.

The Fe2p3/2 core level spectra were also similar after differentsputtering time: 0s, 10s and 3490s, as shown in Fig. 4d. It shows amain feature at a BE of 711.0 ± 0.2 eV, which is characteristic ofFe3+, after peak decomposition analysis [42]. The peak located atthe 709.6.0 ± 0.2 eV which corresponded to Fe2+ was not obviousin the present study. The iron trivalent oxide should correspondto the spinel oxides containing Fe detected by the GIXRD and alsothe possible Fe(OH)3 [40]. Table 2 and Fig. 3 indicated that the con-tent of Ni in the oxide film was excessive, if considering spinelstructure of the oxide. So Fe in oxide film should be present atthe form of Fe(III).

3.4. TEM analysis

3.4.1. Plane-view observation of oxide filmFig. 5a–c shows the bright field images of the big oxide particle

on the MP surface, needle-like oxide on the EP surface and sheet-like oxide on the ground surface respectively. The big oxide particle

P Alloy 690TT after 2 months’ immersion in 1500 mg/L B + 2.3 mg/L Li solution withra of Cr2p3/2; (d) spectra of Fe2p3/2.

Fig. 5. TEM micrographs with electron diffraction patterns of the oxides formed on Alloy 690TT after 2 months’ immersion in 1500 mg/L B + 2.3 mg/L Li solution with1.40 cm3 (STP) O2/kg H2O at 325 �C: (a) big oxide particles formed on the MP surfaces; (b) needle-like oxides formed on the EP surfaces; (c) sheet-like oxides formed on theground surfaces.

Table 3Chemical compositions of oxides shown in Fig. 5 by EDS analysis of TEM.

Ni(at.%)

Fe(at.%)

Cr(at.%)

Big oxide particles formed on the MP surfaces 46.9 49.8 3.3Needle-like oxides formed on the EP surfaces 74.7 23.0 2.3Sheet-like oxides formed on the ground

surfaces70.8 25.9 3.3

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and sheet-like oxide were about several microns in size and thediameter of the needle-like oxide was about 70 nm. SupplementalTEM–EDS microanalyses of the big oxide particle, needle-like andsheet-like oxides indicated differences in chemical composition,as listed in Table 3. It was obvious that the big oxide particle wasrich in Fe and Ni but depleted in Cr, and the contents of Fe andNi were similar. This was consistent with the results in Table 2.The needle-like and sheet-like oxides were also rich in Ni and Febut depleted in Cr, and the content of Ni was higher than that ofFe. The electron diffraction patterns of the big oxide particles, nee-dle-like and sheet-like oxides (inserted in Fig. 5a–c), combinedwith their stoichiometry by the EDS microanalysis indicated thatall these oxides had spinel-type crystal structure (NixCryFezO4),which was consistent with the results of GIXRD that the outer layerof the oxide film was composed of spinel oxides. As the atomic per-cent of Ni, Cr, and Fe were not stoichiometric, it was difficult to de-duce the chemical formula of the spinel oxide.

3.4.2. Cross-sectional observation of oxide filmThe cross-sectional morphologies of oxides films on the surface

ground to 400 grit (Fig. 6) and EP surface (Fig. 7) indicated thatboth of the oxide films had three-layer structure. The big polyhe-dral oxide particles scattered at the outmost of the oxide filmswas the first layer (Figs. 6a and 7a). The loose sheet-like oxideson the ground surface (Fig. 6a and b) or needle-like oxides on theEP surface (Fig. 7a and b) located at the intermediate of the oxidefilms was the second layer, and the irregular cellular oxide of100 nm to several hundred nanometers in size located at innerlayer of the oxide films was the third layer (Figs. 6c and 7c).

Figs. 6 and 7 clearly show that big oxides particles seem to beembedded on the sheet-like and needle-like oxides and were re-deposited. It should be noticed that there was a very clear interfacebetween the intermediate layer and inner layer as shown in

Figs. 6b and 7b, from which the needle-like or sheet -like oxidesand the inner cellular oxides started to grow. The total thicknessof the observed oxide films grown on ground and EP surfaces wereabout 8–9 lm. The thicknesses of the inner layers of the oxidefilms on the ground and EP surfaces were about 1400 nm, andthe inner layers were also very porous. There were some obviousholes in the inner layer just below the interface.

Fig. 8 shows the EDS area mapping of the oxide films. The map-ping areas are shown in Fig. 8a and b by the big rectangles. Thesmall rectangles were used to correct the drift during analysis.The element distributions of Ni, Cr, Fe and O in the oxides filmson the ground and EP surfaces were similar. The outmost layerwas Fe and Ni rich and the intermediate layer was Ni rich, whichwere similar with the results in Tables 2 and 3. For the inner layer,Figs. 6c and 7c indicated the contrast between the center part andthe boundary of irregular cellular oxides was different. EDS map-ping verified that the center of cellular oxides was Ni rich whereasCr was concentrated at the boundaries of cellular oxides. The dif-fraction pattern shown in Fig. 9a indicated that the center of cellu-lar oxide had NiO crystal structure. It was spinel-type diffractionring (such as NiCr2O4) for the oxides at the boundary of cellularoxide (Fig. 9b), indicating the oxides at the boundaries were com-posed of nanostructured spinel oxides.

Fig. 10a shows a preferentially oxidized ribbon at the advancingedge of the oxide film on EP Alloy 690TT after 2 months’ immer-sion. Part of this ribbon was bombarded away by gallium ions dur-ing sample preparation. Its width was about 160–360 nm and theobserved length was about 1.67 lm. The convergent beam diffrac-tion was performed on the matrix located at the two sides of thisribbon under the same tilting condition, as shown in Fig. 10b andc. Different Kikuchi diffraction patterns from the two sides indi-cated that this ribbon was a preferentially oxidized grain bound-ary. Therefore, intergranular attack (IGA) occurred on Alloy690TT in the present study and the IGA rate was bigger than10.16 lm/years. If considering the thickness of the surface oxidefilm, the IGA rate was bigger than 18.67 lm/years. The EDS linescan profiles of this oxidized grain boundary shown in Fig. 10dindicated that this oxidized grain boundary was Ni rich but Cr de-pleted severely, compared with the composition of matrix. Similarwith the composition of the irregular cellular oxides in the innerlayer, Cr mainly concentrated at the boundary of this grain bound-ary. The diffraction ring of the oxides within this ribbon (Fig. 10e)indicated that this oxidized grain boundary was composed ofnanostructured nickel oxides.

Fig. 6. (a) Cross-section of oxide film grown on Alloy 690TT ground to 400 grit after 2 months’ immersion in 1500 mg/L B + 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/kg H2Oat 325 �C; (b) the enlargement of sheet-like oxides/irregular cellular oxides interface; (c) the enlargement of the advancing edge of the oxide films.

Fig. 7. (a) Cross-section of oxide film grown on EP Alloy 690TT after 2 months’ immersion in 1500 mg/L B + 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/kg H2O at 325 �C; (b)the enlargement of needle-like oxides/irregular cellular oxides interface; (c) the enlargement of the advancing edge of the oxide films.

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Fig. 8. TEM image of the oxide film and EDS mapping of the oxide film grown on Alloy 690TT after 2 months’ immersion in 1500 mg/L B + 2.3 mg/L Li solution with 1.40 cm3

(STP) O2/kg H2O at 325 �C: (a) STEM image of oxide film grown on Alloy 690TT ground to 400grit; (b) STEM image of oxide film grown on EP Alloy 690TT; (1) O mapping; (2)Cr mapping; (3) Ni mapping; (4) Fe mapping.

Fig. 9. Diffraction pattern of the oxides at: (a) the center of irregular cellular oxides;(b) the boundaries of the irregular cellular oxides.

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4. Discussions

4.1. Growth of triple-layered surface oxide films on Alloy 690TT

The oxide film observed in this study had a triple-layer struc-ture, as shown in Figs. 6 and 7. The layered structure and oxidephase distribution of the oxide films analyzed by GIXRD and TEMwas also consistent with each other. After observing the structureof oxide film carefully and considering the possible growth mech-anism of every layer in the present study, it could be excluded an

assumption first that the needle-like and sheet-like oxides weredeposited from the solution, as it was not observed on the othersamples (such as pure Cr and 316L stainless steels) in the sameimmersion experiment. It was sheet-like oxide on the groundand MP samples (Fig. 1a–b) whereas it was needle-like oxide onthe EP sample (Fig. 1d). The growth of the intermediate layer mustbe related with the microstructures of the surfaces. The inner layerwas continuous and the cellular oxides were grown up in the ma-trix (Figs. 6c and 7c). Moreover, the intermediate layer started togrow outward and the inner layer stared to grow inward fromthe clear interface (Figs. 6b and 7b). So it was concluded that thisinterface between the intermediate and inner layer should be thenominally original surface of the sample. The intermediate layerwith sheet-like or needle-like oxides was formed by external cor-rosion; and the inner layer with cellular oxides was formed byinternal corrosion. The Pt-marked Alloy 690TT sample was also im-mersed in the same solution for 30 days and the cross section ofoxide film will be observed to further verify the above growthmechanism. Fig. 11 schematically shows the triple-layer structureof oxide films on the ground and polished surfaces, where thehydroxides identified by XPS were also included.

4.1.1. Growth of inner layerThere were no data available for the diffusion coefficient of oxy-

gen along grain boundaries for Alloy 690TT. Considering Alloy690TT was nickel-based, Rebak et al. suggested that the data forNi could be taken. It was about 8.25 � l0�11 cm2/s at 330 �C forthe grain boundary diffusion coefficient of oxygen [49]. Marchettiet al. found grain boundary diffusion coefficient of oxygen in

Fig. 10. (a) Preferentially oxidized ribbon on oxide film grown on EP Alloy 690TT after 2 months’ immersion in 1500 mg/L B and 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/kg H2O at 325 �C and 15.6 MPa; (b), (c), convergent beam diffraction patterns of the matrix located at the two sides of this ribbon; (d) EDS line scan profiles of this ribbonshown in (a) by the black line; (e) diffraction patter of the oxides in this ribbon.

Z. Zhang et al. / Corrosion Science 53 (2011) 3623–3635 3631

chromate like oxide was estimated to be in the range 2 � 10�18 �1 � 10�17 cm2/s and the bulk diffusion coefficient was about8 � 10�22 cm2/s for the oxide film formation on Alloy 690 in thehydrogenated primary water [50]. In the present study, the pene-tration rate of oxygen along the grain boundary of Alloy 690TTwas about 10.16 lm/years, which was also very fast (Fig. 10).

The potential-pH diagrams of Ni–H2O, Cr–H2O and Fe-H2O at300 �C indicated that the stable forms of Ni, Cr and Fe in the hightemperature water with 1.40 cm3 (STP) O2/kg H2O were Ni oxides,dissoluble CrO4

2� and FeO42� respectively [39]. TEM–EDS mapping

in Fig. 8 shows that the whole oxide film was depleted in Cr se-verely and the content of Fe in the outmost and intermediate layerswas very high. So most of Cr in the matrix preferentially dissolvedinto the solution and part of Fe was stabilized in the surface oxides.The growth of this layer is mainly related with the in situ oxidationof Ni in the matrix after the leaching of Cr and Fe, following the in-ward diffusion of the oxidants along the shortcuts, such as grainboundaries and dislocations. The inner layers for all surface sta-tuses were cellular Ni-rich oxide wrapped by nanostructured Cr-rich oxide. It seems the Cr-rich oxides at the boundaries stoppedthe cellular nickel oxides growing. The proposed growth mecha-nism of these cellular oxides was explained as following.

At first, the oxidants could diffuse into the matrix along theshortcuts and most of the Cr and Fe in the matrix would diffuseto the alloy surface. Only trace level of Cr was oxidized and thenstabilized in the matrix. So nickel oxides would nucleate in the in-ner layer. The surrounding nickel and oxygen atoms would diffuseto the nucleation sites and the nickel oxide nucleus could grow up.

Next, it seemed that limited Cr oxides would concentrate at theperiphery of the nickel oxides. With the development of the oxida-tion progress, nickel oxides grew bigger and most of Cr in the oxi-dized matrix dissolved into the solution. Limited Cr oxidescontinued to concentrate at the periphery of the nickel oxides.When nickel atoms around the growing nickel oxides were con-sumed or the nickel oxides were enclosed completely by Cr oxideswhich could play the role of barrier layer well, the diffusion ofnickel atoms and oxidants into the growing nickel oxide particleswas impeded. As a result, the growth of irregular cellular nickeloxide was stopped and a nickel oxide enclosed by Cr oxides wasformed. Then a new nickel oxide would nucleate and grow follow-ing the same manner.

It had been discussed that only trace level of Cr in the matrixwas stabilized in the inner layer. Those Cr-rich oxides should bestable in the inner layer where there was no penetrated water,such as the advancing edge of the inner layer (Figs. 6c and 7c).However, those Cr-rich oxides might be not stable in the innerlayer just below the interface where the water and oxygen couldreach by diffusion. Figs. 6b and 7b showed that the inner layer justbelow the interface was very loose and there were some obviousholes, indicating the dissolution of some oxides. Kim suggestedthat at high electrochemical corrosion potential associated withoxidizing species, the solubility of Cr was higher than that of Feand Ni since Cr tended to go into solution as chromate (CrO4

2�)by the preferential dissolution of Cr oxides [51]:

ðNi; FeÞCr2O4 þ 1:5O2 þH2O! ðFe2þ;Ni2þÞ þ 2CrO2�4 þ 2Hþ ð1Þ

Fig. 11. Schematic drawing for the growth of oxide films on Alloy 690TT with different surface conditions in 1500 mg/L B + 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/kg H2Oat 325 �C and 15.6 MPa.

3632 Z. Zhang et al. / Corrosion Science 53 (2011) 3623–3635

4.1.2. Growth of intermediate layerThis layer is mainly formed by the external corrosion due to the

outward diffusion of the metallic cations from the matrix along theshortcuts in the matrix [52]. It was sheet-like oxide on the groundand MP surfaces as shown in Fig. 1a–c, and needle-like oxide on EPsurface as shown in Fig. 1d. The diffraction patterns of needle-likeand sheet-like oxides shown in Fig. 5b and c indicated that twokinds of oxides had the spine-type structure. However, Table 3shows that the sheet-like and needle-like oxides were rich in Niand Fe, but depleted in Cr. The ratio of the concentration of Ni toFe in the sheet-like and needle-like oxides was about 2.73 and3.25 respectively. If considering the spinel structure, there wasexcessive Ni in these oxides. Hou et al. studied the corrosion of al-loy 600 with hydrogen charged and uncharged in the high-puritywater with 5.6 cm3 (STP) O2/kg H2O at 288 �C and 8.5 MPa[53,54]. It was also found that the spinel-type needle-like or pyra-mid-like oxides were rich in Ni and the ratio of Ni to Fe was about1.90 and 2.80 for the uncharged and charged samples. Hou sug-gested that the needle-like oxides might be nickel-based [54].

For the ground and EP samples, the surface dislocations of highenergy were the preferential nucleation sites for the oxides. So thespinel oxides would nucleate at these sites quickly and then growup by the association between the oxygen and the outward diffu-sion Ni, Fe and Cr cations, during which the adsorbed oxygenatoms were reduced. The reduced oxygen atoms should be pro-vided by the dissolved oxygen or by the water molecular, as thedetection of hydroxides indicated that water molecular was alsoincorporated into the growth of surface oxide film (Fig. 4). The re-lated reactions mainly occurred at the sample surfaces. The growthof spinel oxides on Alloy 690TT in the high temperature watermight be an electrochemical process and might be different fromthe oxidation in the dry oxygen or wet oxygen where spinel oxideswere formed by the solid state reactions of different oxides.

4.1.3. Growth of the outmost dispersed big polyhedral oxide particlesThe fact that these scattered big oxide particles were embedded

in the intermediate layers but not started to grow from the originalmatrix surface (Figs. 1, 6 and 7) suggested the growth of these

oxide particles agreed well with the dissolution and precipitationmechanism, which had been discussed widely elsewhere [15,30–32,55]. McIntype et al. suggested that the reprecipitation of thedissolved Ni could be negligible [18]. So these oxide particles richin Fe and Ni were believed to grow by recrystallization of iron ionspicked up from the test solution with the outwards diffused nickel(II), Fe (III) and chromium (III) ions from the matrix [30–32].

However, Sennour et al. had observed the metastable solid par-tially hydrated nickel hydroxide (Ni(OH)2�yH2O) in the externallayer of the oxide film developed on Ni–30Cr alloy exposed inthe simulated hydrogenated primary water [40]. Its formationwas verified to result from the precipitation of the neutral aqueouscomplex Ni(OH)2(aq). Combined with the iron hydroxide whose for-mation was similar with that of nickel hydroxide, the nucleationand growth of nickel ferrite oxide particles was occurred by the fol-lowing reaction [40]:

ð2þ zÞFeðOHÞ2 þ ð1� zÞNiðOHÞ2$ Nið1�zÞFeð2þzÞO4 þ 2H2OþH2ðaqÞ ð2Þ

The concentrations of Fe and Ni cations in the solution could af-fect the stochiometry of nickel ferrite.

EDS analysis indicated the content of Ni in the big oxide parti-cles were very high, as listed in Table 2 and 3. There must be exces-sive Ni, if considering the spinel-type structure (such as NiFe2O4).According the above discussions, the excessive Ni in the big oxideparticles might be caused by the reprecipitation of Ni from thesolution, which should not be negligible. Trace level of Cr in theoxide particles and chromium hydroxides detected by the XPSmight also be precipitated from the solution, which had been dis-cussed by Sun et al. [36]. As a result, the reprecipitation of Fe, Niand Cr occurred for Alloy 690TT and the outmost layer with oxideparticles and hydroxides were formed in the present oxygenatedprimary water (Fig. 11).

The used test loop including autoclave is made of 316L stainlesssteel, which could also release metallic ions to the solution andthen affect the growth of big oxide particles. The metallic ionspick-up phenomenon is believed to be responsible to the nonstoi-chiometric spinel structure of the big oxide particles.

Z. Zhang et al. / Corrosion Science 53 (2011) 3623–3635 3633

4.2. Effect of surface status on oxidation of Alloy 690TT

Different surface finish would cause different superficial cold-worked layer on Alloy 690TT [56]. The cold-worked layers on theground and MP surfaces had refined grains, and no obvious coldworked layer existed on the EP surface. The dislocation and sub-grain boundary densities on the ground surfaces were much higherthan those on the EP surface. The main defects at the near-surfacelayer of the ground samples were mainly parallel dislocation lines.

4.2.1. Effect of surface status on oxides morphologiesComparing the oxide films grown on ground and EP samples,

the main differences were the morphologies of the intermediateoxide layer. The intermediate layer of the oxide film on EP surfacewas needle-like oxides (Figs. 1d and 7), which was consistent withthe observation on the surfaces of Alloy 600 and 5 wt% Cr alloyafter immersion in high temperature hydrogenated water[33,57]. Fig. 12a shows the TEM plane-views of a single needle-likeoxide. The contrast between the boundary and the center of

Fig. 12. (a) TEM contrast observation of needle-like oxides grown on EP Alloy 690TT afterkg H2O at 325 �C and 15.6 MPa; (b) cross-sectional EDS line scan profiles of a needle-lik

Fig. 13. Size and density comparison of the big oxide particles grown on Alloy 690TT wiB + 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/kg H2O at 325 �C and 15.6 MPa: (a) groun

needle-like oxide was different and a very clear ribbon of 15–20 nm in width was found at the center of the needle-like oxide.The EDS line scan profiles indicated that the intensities of the ele-ments Ni, Cr and Fe at the center were much higher than those atthe both edges (Fig. 12b), indicating the element distribution in thecross-section of the needle-like oxides was non-homogeneous. Thepresent results were compatible with the existence of centred tun-nel for atom diffusion in the needle-like oxide [58]. Voss et al. hadconfirmed the existence of a central tunnel within the a-Fe2O3

whisker on iron from 600 to 800 �C by TEM observation [58]. Oxidewhiskers grown in oxygen containing water vapor were extremelylong and thin and water vapor promoted the growth along the axis[59]. So the nucleation and growth of the needle-like oxides on EPsurface with dislocations of low density in high temperature watercould be explained well by the growth mechanism of oxide whis-kers, which had been discussed widely [58–65]. The formation ofneedle-like oxides was mainly related with the fast surface diffu-sion of cations along a tunnel centred on the core of a screw dislo-cation or a bundle of screw dislocations [58–65]. The local flux

2 months’ immersion in 1500 mg/L B + 2.3 mg/L Li solution with 1.40 cm3 (STP) O2/e oxide.

th different surface states after immersion after 2 months’ immersion in 1500 mg/Ld to 400 grit; (b) ground to 1500 grit; (c) mechanical polished; (d) electropolished.

3634 Z. Zhang et al. / Corrosion Science 53 (2011) 3623–3635

associated with the surface diffusion along with a central whiskercore could be several orders of magnitude larger than the flux asso-ciated with the volume diffusion [59]. A screw dislocation emerg-ing on the surface of a crystal provided the reproducible surfacestep needed for the lattice growth of needle-like oxides [59]. Sce-nini et al. observed the oxide nodules on the electropolished alloy600 and 690 after exposure in a steam-hydrogen mixture [66]. Foralloy 600, to form a feature 200 nm high in 65 h required a diffusiv-ity of at least the order of 10�19 m2/s, which was more than 4 or-ders of magnitude higher than the lattice elf-diffusivity of Ni at480 �C which was 7 � 10�24 m2/s. The growth of nodules was alsodominated by fast surface and interface diffusion [66].

Compared to the EP surface, the ground surface had a cold-worklayer with dislocation lines of high density and more grain bound-aries due to grain refinement, which could provide more shortcutsfor diffusion and more preferential sites for oxide nucleation [67].As a result, more Ni, Cr and Fe atoms could be transported quicklyto the sample surface. The spinel oxides would preferentiallynucleate along those surface dislocation lines. In oxidizing environ-ment, the oxidants were enough and metallic atoms could associ-ate with the adsorbed oxidants quickly, as soon as metallic atomswere transported to the sample surface. So the lateral diffusion ofthe cations was constrained. The oxides could only grow toward adirection that was vertical to the sample surfaces through a latticegrowth at the original nucleation sites. Then the sheet-like oxideswere formed.

4.2.2. Effect of surface status on the oxidation rateThe nucleation and growth of the external corrosion layer on

the ground surface were much easier those that on the EP surface.However, the oxide films grown in the present study was loose anddepleted in Cr. It had been discussed that the nucleation andgrowth of outmost big oxide particles would be suppressed whena protective layer of chromium-rich inner oxide was established[55]. However, the size and density of big oxide particles on differ-ent surfaces were similar and were also very big (Fig. 13); indicat-ing oxide films were lack of protectiveness. So the growth of oxidefilms on the ground and EP surfaces would be unrestrained. After2 months’ immersion, the thicknesses of the inner oxide layerson different surface status are about 1400 nm, which was muchthicker than the superficial cold-worked layers (285–470 nm)[56]. So the effect of surface statuses on the oxidation rate becomeobscure and the internal oxidation rates on the ground and EP sam-ples were about 8.52 lm/years.

4.3. The IGA sensitivity of Alloy 690TT in oxygenated primary water

Yu and Yao had investigated the relationship between the resis-tance to IGA and intergranular stress corrosion cracking (IGSCC)and the chromium depletion of Alloy 690, and indicated that high-er equilibrium chromium concentration at the intergranular car-bide/matrix interface results in good resistance to IGA and IGSCC[68]. However, Cr in the matrix was thermodynamically unstableand the oxide film was lack of protectiveness when oxygen wasdissolved into the primary water. The preferential dissolutionand oxidation of Cr at the grain boundaries would occur. Panteret al. and Ovanessian et al. had discussed that the growth of theoxides layers developing in primary water could create defectsby itself in the material [6,69]. The selective oxidation of chromiumcould produce vacancies close to the metal-oxide interface. Drivenby concentration gradient these vacancies would diffuse towardsthe bulk of the material, essentially by the grain boundaries. Thefirst effect of this migration of vacancies was an acceleration ofthe diffusion of chromium in the opposite direction acceleratingthe chromium depletion. The second effect of this migration ofvacancies was that a transport of the interstitial species such as

oxygen was expected by a binding effect between the vacanciesand these atoms of small sizes [6,69]. As a result, the inward pen-etration of oxygen at the grain boundaries was accelerated for Al-loy 690TT and a clear IGA was found at the advancing edge of theoxide film on EP surface after immersion for 2 months. Based onthis TEM sample, the IGA rate was bigger than 10.16 lm/years(Fig. 10).

Therefore, the dissolved oxygen in the high temperature pri-mary water not only caused the formation of un-protective inneroxide layer but also increased the sensitivity of IGSCC and IGA ofAlloy 690TT a lot. Dissolved oxygen should be forbidden in primaryside during start-up and operation of PWRs when Alloy 690TT wasused as the steam generator tubing material.

5. Conclusions

The oxide films grown on Alloy 690TT samples with differentsurface conditions in the oxygenated primary water were system-atically characterized. The following conclusions could bewithdrawn:

1) Triple-layer oxide films were formed on Alloy 690TT. Theoutmost big polyhedral oxide was formed by reprecipitation,the intermediate layer was formed by external corrosion andthe inner layer was formed by internal corrosion.

2) The porous inner layer on both ground and electropolished sur-faces were Ni-rich cellular oxide wrapped by nano-structuredCr-rich oxide. The intermediate layers were loose Ni-richsheet-like oxides on the ground surface and needle-like oxideson the electropolished surface. The outmost layers were dis-persed big polyhedral oxide particles and rich in Ni and Fe.

3) Surface status only affected the intermediate layer oxidemorphology, but did not affect internal corrosion rate. Theinner corrosion rate for both surfaces was about 8.52 lm/years. The oxide films grown on different surfaces couldnot protect the matrix from corrosion in the present oxygen-ated primary water.

4) Dissolved oxygen increased the IGA sensitivity of Alloy690TT and should be forbidden in primary side duringstart-up and operation when Alloy 690TT was used as thesteam generator tubing material.

Acknowledgments

This work is supported by the Special Funds for the Major StateBasic Research Projects G2011CB610502 and Nature Science Foun-dation 51025104. The authors gratefully acknowledge EPRI for pro-viding commercial bulk Alloy 690TT and tubing, also thankShanghai Synchrotron Radiation Facility (SSRF) and National Syn-chrotron Radiation Laboratory of University of Science and tech-nology China (NSRL) for using GIXRD.

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