- Email: [email protected]
Precipitation of cotunnite-type zirconia induced by deuterium absorption in Zr-4 alloy
Scripta Materialia 143 (2018) 1–4
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular article
Precipitation of cotunnite-type zirconia induced by deuterium absorption in Zr-4 alloy Cheng Zhang, Yun Yang, Yin Zhang, Jingru Liu, Li You, Xiping Song ⁎ State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 11 May 2017 Received in revised form 20 July 2017 Accepted 1 August 2017 Available online xxxx Keywords: Cotunnite-type zirconia Deuterium absorption TEM Zr-4 alloy
a b s t r a c t Cotunnite-type zirconia was first observed to precipitate from Zr-4 alloy after deuterium absorption. The acicular cotunnite-type zirconia (OII-ZrO2) precipitated at the interface of δ deuterides, followed by an orientation relationship of ½011OII ==½011δ and ð100ÞOII ==ð111Þδ with δ deuterides. High-resolution transmission electron microscopy observations indicated that some steps existed at the δ/ZrO2 interface, together with the appearance of dislocations in the δ deuteride side. A model of the formation of cotunnite-type ZrO2 at the interface of δ deuterides was proposed. The effects of volume strain and crystal structure on the precipitation of cotunnite-type ZrO2 were also analyzed. © 2017 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Zirconium alloys have been widely used as cladding materials in nuclear reactors because of their low thermal neutron absorption section, good mechanical properties, and corrosion resistance [1]. During reactor operation, zirconium is oxidized in an aqueous environment by the reaction Zr + 2H2O → ZrO2 + 4H. A fraction of the hydrogen produced in the oxidation reaction is absorbed by the zirconium alloy matrix and precipitates as hydrides when the hydrogen content exceeds its solubility. The precipitation of hydrides leads to a considerable degradation in the mechanical properties of the zirconium alloys [2–4], which are highly dependent on the morphology [5], distribution [6], and orientation [7,8] of these hydrides. In addition to the oxidation of zirconium alloys, the hydrogenation of zirconium alloys is another important issue with regard to the accelerated corrosion of zircaloy cladding at high burn-up [9–11]. Studies have shown that the precipitation of hydrides affects the transition of oxides at the zirconia/Zr interface [12–14]. It has been reported that the proportion of tetragonal zirconia, which acts as a dense corrosion resistance layer, is less in pre-hydride Zircaloy-4 (Zr-4) and Zr–1.5Nb alloys [15]. Zr3O has been reported to form between oxide film and hydride rim [13] and is regarded as the intermediate product of the transformation from δ-hydride to oxide. Another phase of ZrO and a disordered amorphous zirconium oxide layer have also been observed at the zirconia/Zr interface in pre-hydride zirconium alloys [15,16]. These studies have mainly focused on the effect of hydrides on the transition of oxides at the oxide/matrix interface. However, within the Zr matrix,
⁎ Corresponding author. E-mail address: [email protected] (X. Song).
http://dx.doi.org/10.1016/j.scriptamat.2017.08.001 1359-6462/© 2017 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
hydrides and oxygen atoms, which are about 1200 ppm in the Zr-4 alloy, also exist. To date, no attention has been paid to the formation of oxides in the Zr matrix after hydrogen absorption. In this article, the microstructures of Zr-4 alloy matrix were studied by transmission electron microscopy (TEM) after deuterium absorption at 1173 K. We used deuterium rather than hydrogen for convenience. Deuterium is an isotope of hydrogen; therefore, deuterium absorption is almost the same as hydrogen absorption. The composition of the as-received Zr-4 alloy, which was supplied by the State Nuclear Bao Tai Zirconium Industry Company, is listed in Table 1. The as-received Zr-4 alloy was annealed at 1073 K for 5 h in vacuum to obtain equiaxed α-Zr and sectioned to samples with a dimension of 4 × 4 × 8 mm for deuterium absorption. Deuterium absorption was performed in a Sievert's apparatus using high purity deuterium (99.999%) at 1173 K and 3 bar, followed by furnace cooling to room temperature at a cooling rate of approximately 5 °C/min. The phase structures were identified by XRD using Cu K-alpha radiation with a step of 0.02° and 2 s/step. The microstructure was characterized by SEM using back-scattered electron. The microstructure and phase structure were investigated using a FEI Tecnai F30 at an accelerating voltage of 300 kV. The foil samples for TEM observation were prepared by twin-
Table 1 Chemical composition of zircolay-4 used in this study. Element
Sn
Amount of weight (wt%)
1.2–1.7 0.18–0.24 0.07–0.13 0.125 0.004 Balance
Fe
Cr
O
H
Zr
2
C. Zhang et al. / Scripta Materialia 143 (2018) 1–4
Fig. 1. (a) Comparison of the XRD patterns of the deuterated (at 1173 K and 3 bar) and as-received samples and (b) backscatter electron image of the sample deuterated at 1173 K and 3 bar.
jet electropolishing using an electrolyte containing 10% HClO4 and 90% ethanol at 243 K. STEM-EELS analysis was performed on a JEOL ARM200F microscope equipped with Gatan image filter (GIF) quantum
965. The atomic composition was analyzed using the Gatan Digital Micrograph EELS module, using the Zr-M4,5 and O-K edges, by the k-factor approach.
Fig. 2. (a) TEM image of ZrO2 and (b) corresponding diffraction patterns from different zone axes; (c) HADDF image of ZrO2 and (d) corresponding EELS spectrum of spots marked in (c); (e) composition profile along the line in (c).
C. Zhang et al. / Scripta Materialia 143 (2018) 1–4
3
Fig. 3. (a) Bright-field image of ZrO2 and δ deuteride platelets, (b) the diffraction pattern of the circled area marked in (a), and (c) HRTEM image of the δ/ZrO2 interface.
Fig. 1 shows the XRD pattern and backscatter electron image of samples deuterated at 1173 K and 3 bar. The XRD results in Fig. 1a show strong peaks of δ deuteride (δ-ZrD 1.66 ) and α-Zr, together with some weak peaks of γ deuteride (γ-ZrD). No zirconia phase was found before and after the deuterium absorption. The BSE image in Fig. 1b shows the morphology of the deuterides (dark area) and α-Zr (bright area). Fig. 2a and b shows the TEM image and the corresponding diffraction patterns with different zone axes. Both the δ deuteride grains and platelets can be observed in Fig. 2a. Between two δ deuteride platelets, there is an acicular precipitate, which is marked by black arrows in Fig. 2a, having a width of about 100 nm. During the TEM observation, the width of the acicular precipitate remained the same during the sample tilting, thus indicating that the acicular precipitate has a bulk feature rather than a surface feature. From the diffraction patterns obtained from different zone axes, as shown in Fig. 2b, the acicular precipitate is indexed to be cotunnite-type ZrO2 (termed OII-ZrO2) with a space group of Pnam. Furthermore, its lattice parameters are measured to be a = 0.548 nm, b = 0.604 nm, and c = 0.335 nm, which are in good agreement with the PDF data (#49-1746, a = 0.5593 nm, b = 0.6484 nm, and c = 0.3333 nm). Fig. 2c and d shows the HADDF image and corresponding EELS spectrum of the acicular precipitate, respectively. The HADDF image shows that the precipitate has a dark contrast. Furthermore, the EELS spectrum shows strong signals of Zr M4,5 and O K edges at 180 and 532 eV, respectively. The O/Zr atom ratio of the precipitate was calculated to be 2.04, which obeys the stoichiometric ratio of OII-ZrO2. Fig. 2e shows the composition profile across the precipitate of OII-ZrO2: the oxygen concentration increases distinctly in the OII-ZrO2 precipitate. Fig. 3a and b shows the morphology of the precipitated OII-ZrO2 phase and the orientation relationship between the δ-deuteride and OII-ZrO2 phase. It can be seen that the acicular OII-ZrO2 precipitates at the interface between the δ-deuteride platelets. The diffraction patterns at the δ/ZrO2 interface, marked as a circle in Fig. 3a, are shown in Fig. 3b. They follow an orientation relationship of ½011OII ==½011δ , ð100ÞOII == ð111Þδ . Fig. 3c shows the HRTEM image of the δ/ZrO2 interface. It can be seen that the δ/ZrO2 interface is parallel with their habit planes,
ð111Þδ and (100)OII. In addition, some steps exist at the δ/ZrO2 interface. Dislocations are observed around these steps on the δ deuteride side. The formation of these steps and dislocations may be caused by the misfit between ð111Þδ and (100)OII because of their different interplanar spacings. From the ICDD crystallography data and the as-determined orientation relationship, a model of the transformation from δ deuteride to OIIZrO2 phase was proposed, as shown in Fig. 4. It shows that the ð111Þδ plane is parallel to the (100)OII plane. It can be seen that the ð111Þδ
Fig. 4. Model illustrating the transformation depending on the orientation relationship of ½011δ ==½011ZrO2 and ð111Þδ ==ð100ÞZrO2 . The green circles represent Zr atoms, gray circles represent D atoms, and red circles represent O atoms. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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C. Zhang et al. / Scripta Materialia 143 (2018) 1–4
Table 2 Calculated volume expansion and lattice misfit in different ZrO2 phases. Phase
Structure
δ-ZrH1.66 Fm3m P21/c P42/nmc Fm3m Pbcm Pnam
m-ZrO2 t-ZrO2 c-ZrO2 OI-ZrO2 OII-ZrO2
V1/nm3
V2/nm3
Ω/%
Matched planesa
Interface misfit
PDF#
/
0.02725
/
/
/
65–6972
0.03650 0.03335 0.03371 0.03472 0.03022
/ / / / /
33.9 22.4 23.7 27.4 10.9
(111)δ//(111)m (111)δ//(011)t (111)δ//(111)c (111)δ//(002)OI (111)δ//(200)OII
3.9% 6.7% 6.9% 5.1% 2.1%
65–2357 50–1089 49–1642 37–1413 49–1746
a Matched planes of these zirconia are the planes with the closest interplanar distance with (111)δ planes according to the standard PDF cards.
plane could transform into the (100)OII plane by increasing the atomic distance along the [011]δ direction from 0.337 to 0.365 nm and decreasing the atomic distance along the [101]δ direction from 0.337 to 0.334 nm and simultaneously enlarging the angle between [011]δ and [101]δ from 60° to 62.8°. Furthermore, the diffusion of deuterium and oxygen is also required. The δ deuteride can then be transformed into the OII-ZrO2 phase. There are several kinds of zirconia. Under atmospheric pressure, stable zirconia is monoclinic m-ZrO2 (P21/c) up to 1440 K, then tetragonal t-ZrO2 (P42/nmc) from 1440 to 2950 K, and finally cubic c-ZrO2 (Fm3m) as the temperature becomes higher than 2950 K [17]. Two orthorhombic ZrO2, denoted as OI and OII, respectively, have been reported to form under higher pressure conditions [18,19]. OI-ZrO2 with a space group of Pbcm has been observed at pressures from 3.5 to 13 GPa; however, it can also be found in doped zirconia, films, and nanoparticles at atmospheric pressure [20–22]. OII structure ZrO2 (space group Pnam) is reported to be stable only at pressures higher than 12.5 GPa [17,18]. However, in this study, OII-ZrO2 phase was found in deuterated samples at atmospheric pressure. The formed OII-ZrO2 phase is not likely to be the surface artifact introduced by foil preparation because it is an acicular precipitate rather than an oxide film. From the observed OII-ZrO2 morphology and the orientation relationship between OII-ZrO2 and δ deuteride, as shown in Fig. 3, it is reasonable to consider that the formation of OII-ZrO2 phase depends on the formation of δ deuteride. During deuterium absorption, dissolved oxygen is expelled and aggregates on the interface of δ deuterides because of the growth of δ deuterides, which have a lower solubility of oxygen than zirconium matrix [23]. This provides a suitable oxygen atom concentration for the formation of zirconia. Furthermore, the formation of OII-ZrO2 may also be related to the volume strain energy. To evaluate the volume strain energy during the formation of zirconia, the atom volume strain Ω was calculated by the following equation [24]: Ω = (V1 − V2) / V2 ∗ 100%, where V1 and V2 are the atom volumes of Zr in ZrO2 and δ deuteride cells, respectively. The calculated atom volume V1 and the atom volume strain Ω of different ZrO2 structures are shown in Table 2. This demonstrates that OIIZrO2 has the lowest volume strain among all kinds of ZrO2 phases—10.9% as compared to 33.9%, 22.4%, 23.7%, and 27.4% for the m-ZrO2, t-ZrO2, c-ZrO2, and OI-ZrO2, respectively. The lowest Ω of OIIZrO2 represents the lowest volume strain energy that is preferential to precipitate zirconia from the δ-deuteride phase.
Furthermore, the misfit between the ð111Þδ and the (200)OII habit planes is small. From Table 2, it can be seen clearly that this is the lowest degree of misfit compared with the misfits between other ZrO2 phases and δ deuterides. Furthermore, by a slight readjustment of Zr atoms within the ð111Þδ habit plane, i.e., an increase in the atomic distance along the [011]δ direction and a decrease in the atomic distance along the [011]δ direction, accompanied by a tiny lattice rotation, as shown in Fig. 4, the FCC lattice of δ deuterides can be transformed into the Pnam lattice of OII-ZrO2 phase, as shown by our experimental results. In summary, the acicular cotunnite-type OII-ZrO2 was observed in Zr-4 alloy after deuterium absorption at 1173 K. It precipitated from the interface of δ deuterides, followed by the orientation relationship of ½011OII ==½011δ and ð100ÞOII ==ð111Þδ . Some steps existed at the δ/ ZrO2 interface, together with the appearance of dislocations in the δ deuteride side. The effects of oxygen content, volume strain, and crystal structure play critical roles in the formation of OII-ZrO2. The discovery of OII-ZrO2 phase after the deuterium absorption will be helpful in the understanding of the oxidation and corrosion behavior of zirconium alloys. This work was supported by the National Natural Science Foundation of China (Nos. 21171018 and 51271021), the Beijing Natural Science Foundation (No. 2162025), and the State Key Laboratory for Advanced Metals and Materials (No. 2016-T02). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]
[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
S.J. Zinkle, G.S. Was, Acta Mater. 61 (2013) 735. L.A. Simpson, Metall. Trans. A. 12 (1981) 2113–2124. J.B. Bai, D. Frangois, J. Nucl. Mater. 187 (1992) 186–189. R. Dutton, K. Nuttall, M.P. Puls, L.A. Simpson, Metall. Trans. A. 8 (1977) 1553–1562. S. Arsene, J. Bai, P. Bompard, Metall. Mater. Trans. A 34 (2003) 579–588. W. Qin, N.A.P. Kiran Kumar, J.A. Szpunar, J. Kozinski, Acta Mater. 59 (2011) 7010. K.B. Colas, A.T. Motta, J.D. Almer, M.R. Daymond, M. Kerr, A.D. Banchik, P. Vizcaino, J.R. Santisteban, Acta Mater. 58 (2010) 6575. M.A. Vicente Alvarez, J.R. Santisteban, P. Vizcaino, A.V. Flores, A.D. Banchik, J. Almer, Acta Mater. 60 (2012) 6892. A.M. Garde, Zirconium in the Nuclear Industry: 9th International Symposium, ASTM STP 1132, 1991 566. M. Blat, D. Noel, Zirconium in the Nuclear Industry: 11th International Symposium, ASTM STP 1295, 1996 319. N. Ni, D. Hudson, J. Wei, P. Wang, S. Lozano-Perez, G.D.W. Smith, J.M. Sykes, S.S. Yardley, K.L. Moore, S. Lyon, R. Cottis, M. Preuss, C.R.M. Grovenor, Acta Mater. 60 (2012) 7132. M. Harada, R. Wakamatsu, Zirconium in the Nuclear Industry: 15th International Symposium, ASTM STP 1505, 2009 384. M. Tupin, C. Bisor, P. Bossis, J. Chêne, J.L. Bechade, F. Jomard, Corros. Sci. 98 (2015) 478. B. de Gabory, A.T. Motta, K. Wang, J. Nucl. Mater. 456 (2015) 272. Y. Kim, Y. Jeong, S. Son, J. Nucl. Mater. 444 (2014) 349. T. Kim, J. Kim, K.J. Choi, S.C. Yoo, S. Kim, J.H. Kim, Corros. Sci. 99 (2015) 134. H. Fukui, T. Kunisada, T. Fujisawa, K. Funakoshi, W. Utsumi, T. Irifune, K. Kuroda, T. Kikegawa, O. Ohtaka, Phys. Rev. B 63 (2001) 174108. O. Ohtaka, D. Andrault, P. Bouvier, E. Schultz, M. Mezouar, J. Appl. Crystallogr. 38 (2005) 727. J. Haines, J.M. Léger, S. Hull, J.P. Petitet, A.S. Pereira, C.A. Perottoni, J.A. Jornada, J. Am. Ceram. Soc. 80 (1997) 1910. A.H. Heuer, V. Lanteri, S.C. Farmer, R. Chaim, R.R. Lee, B.W. Kibbel, R.M. Dickerson, J. Mater. Sci. 24 (1989) 124. S. Liu, W. Hu, Y. Zhang, J. Xiang, F. Wen, B. Xu, J. He, D. Yu, Y. Tian, Z. Liu, J. Appl. Crystallogr. 47 (2014) 684. G. Trolliard, R. Benmechta, D. Mercurio, Acta Mater. 55 (2007) 6011. D. Setoyama, S. Yamanaka, J. Alloys Compd. 370 (2004) 144. Y. Yang, X.P. Song, C. Zhang, J. Nucl. Mater. 465 (2015) 97.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular article
Precipitation of cotunnite-type zirconia induced by deuterium absorption in Zr-4 alloy Cheng Zhang, Yun Yang, Yin Zhang, Jingru Liu, Li You, Xiping Song ⁎ State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 11 May 2017 Received in revised form 20 July 2017 Accepted 1 August 2017 Available online xxxx Keywords: Cotunnite-type zirconia Deuterium absorption TEM Zr-4 alloy
a b s t r a c t Cotunnite-type zirconia was first observed to precipitate from Zr-4 alloy after deuterium absorption. The acicular cotunnite-type zirconia (OII-ZrO2) precipitated at the interface of δ deuterides, followed by an orientation relationship of ½011OII ==½011δ and ð100ÞOII ==ð111Þδ with δ deuterides. High-resolution transmission electron microscopy observations indicated that some steps existed at the δ/ZrO2 interface, together with the appearance of dislocations in the δ deuteride side. A model of the formation of cotunnite-type ZrO2 at the interface of δ deuterides was proposed. The effects of volume strain and crystal structure on the precipitation of cotunnite-type ZrO2 were also analyzed. © 2017 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Zirconium alloys have been widely used as cladding materials in nuclear reactors because of their low thermal neutron absorption section, good mechanical properties, and corrosion resistance [1]. During reactor operation, zirconium is oxidized in an aqueous environment by the reaction Zr + 2H2O → ZrO2 + 4H. A fraction of the hydrogen produced in the oxidation reaction is absorbed by the zirconium alloy matrix and precipitates as hydrides when the hydrogen content exceeds its solubility. The precipitation of hydrides leads to a considerable degradation in the mechanical properties of the zirconium alloys [2–4], which are highly dependent on the morphology [5], distribution [6], and orientation [7,8] of these hydrides. In addition to the oxidation of zirconium alloys, the hydrogenation of zirconium alloys is another important issue with regard to the accelerated corrosion of zircaloy cladding at high burn-up [9–11]. Studies have shown that the precipitation of hydrides affects the transition of oxides at the zirconia/Zr interface [12–14]. It has been reported that the proportion of tetragonal zirconia, which acts as a dense corrosion resistance layer, is less in pre-hydride Zircaloy-4 (Zr-4) and Zr–1.5Nb alloys [15]. Zr3O has been reported to form between oxide film and hydride rim [13] and is regarded as the intermediate product of the transformation from δ-hydride to oxide. Another phase of ZrO and a disordered amorphous zirconium oxide layer have also been observed at the zirconia/Zr interface in pre-hydride zirconium alloys [15,16]. These studies have mainly focused on the effect of hydrides on the transition of oxides at the oxide/matrix interface. However, within the Zr matrix,
⁎ Corresponding author. E-mail address: [email protected] (X. Song).
http://dx.doi.org/10.1016/j.scriptamat.2017.08.001 1359-6462/© 2017 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
hydrides and oxygen atoms, which are about 1200 ppm in the Zr-4 alloy, also exist. To date, no attention has been paid to the formation of oxides in the Zr matrix after hydrogen absorption. In this article, the microstructures of Zr-4 alloy matrix were studied by transmission electron microscopy (TEM) after deuterium absorption at 1173 K. We used deuterium rather than hydrogen for convenience. Deuterium is an isotope of hydrogen; therefore, deuterium absorption is almost the same as hydrogen absorption. The composition of the as-received Zr-4 alloy, which was supplied by the State Nuclear Bao Tai Zirconium Industry Company, is listed in Table 1. The as-received Zr-4 alloy was annealed at 1073 K for 5 h in vacuum to obtain equiaxed α-Zr and sectioned to samples with a dimension of 4 × 4 × 8 mm for deuterium absorption. Deuterium absorption was performed in a Sievert's apparatus using high purity deuterium (99.999%) at 1173 K and 3 bar, followed by furnace cooling to room temperature at a cooling rate of approximately 5 °C/min. The phase structures were identified by XRD using Cu K-alpha radiation with a step of 0.02° and 2 s/step. The microstructure was characterized by SEM using back-scattered electron. The microstructure and phase structure were investigated using a FEI Tecnai F30 at an accelerating voltage of 300 kV. The foil samples for TEM observation were prepared by twin-
Table 1 Chemical composition of zircolay-4 used in this study. Element
Sn
Amount of weight (wt%)
1.2–1.7 0.18–0.24 0.07–0.13 0.125 0.004 Balance
Fe
Cr
O
H
Zr
2
C. Zhang et al. / Scripta Materialia 143 (2018) 1–4
Fig. 1. (a) Comparison of the XRD patterns of the deuterated (at 1173 K and 3 bar) and as-received samples and (b) backscatter electron image of the sample deuterated at 1173 K and 3 bar.
jet electropolishing using an electrolyte containing 10% HClO4 and 90% ethanol at 243 K. STEM-EELS analysis was performed on a JEOL ARM200F microscope equipped with Gatan image filter (GIF) quantum
965. The atomic composition was analyzed using the Gatan Digital Micrograph EELS module, using the Zr-M4,5 and O-K edges, by the k-factor approach.
Fig. 2. (a) TEM image of ZrO2 and (b) corresponding diffraction patterns from different zone axes; (c) HADDF image of ZrO2 and (d) corresponding EELS spectrum of spots marked in (c); (e) composition profile along the line in (c).
C. Zhang et al. / Scripta Materialia 143 (2018) 1–4
3
Fig. 3. (a) Bright-field image of ZrO2 and δ deuteride platelets, (b) the diffraction pattern of the circled area marked in (a), and (c) HRTEM image of the δ/ZrO2 interface.
Fig. 1 shows the XRD pattern and backscatter electron image of samples deuterated at 1173 K and 3 bar. The XRD results in Fig. 1a show strong peaks of δ deuteride (δ-ZrD 1.66 ) and α-Zr, together with some weak peaks of γ deuteride (γ-ZrD). No zirconia phase was found before and after the deuterium absorption. The BSE image in Fig. 1b shows the morphology of the deuterides (dark area) and α-Zr (bright area). Fig. 2a and b shows the TEM image and the corresponding diffraction patterns with different zone axes. Both the δ deuteride grains and platelets can be observed in Fig. 2a. Between two δ deuteride platelets, there is an acicular precipitate, which is marked by black arrows in Fig. 2a, having a width of about 100 nm. During the TEM observation, the width of the acicular precipitate remained the same during the sample tilting, thus indicating that the acicular precipitate has a bulk feature rather than a surface feature. From the diffraction patterns obtained from different zone axes, as shown in Fig. 2b, the acicular precipitate is indexed to be cotunnite-type ZrO2 (termed OII-ZrO2) with a space group of Pnam. Furthermore, its lattice parameters are measured to be a = 0.548 nm, b = 0.604 nm, and c = 0.335 nm, which are in good agreement with the PDF data (#49-1746, a = 0.5593 nm, b = 0.6484 nm, and c = 0.3333 nm). Fig. 2c and d shows the HADDF image and corresponding EELS spectrum of the acicular precipitate, respectively. The HADDF image shows that the precipitate has a dark contrast. Furthermore, the EELS spectrum shows strong signals of Zr M4,5 and O K edges at 180 and 532 eV, respectively. The O/Zr atom ratio of the precipitate was calculated to be 2.04, which obeys the stoichiometric ratio of OII-ZrO2. Fig. 2e shows the composition profile across the precipitate of OII-ZrO2: the oxygen concentration increases distinctly in the OII-ZrO2 precipitate. Fig. 3a and b shows the morphology of the precipitated OII-ZrO2 phase and the orientation relationship between the δ-deuteride and OII-ZrO2 phase. It can be seen that the acicular OII-ZrO2 precipitates at the interface between the δ-deuteride platelets. The diffraction patterns at the δ/ZrO2 interface, marked as a circle in Fig. 3a, are shown in Fig. 3b. They follow an orientation relationship of ½011OII ==½011δ , ð100ÞOII == ð111Þδ . Fig. 3c shows the HRTEM image of the δ/ZrO2 interface. It can be seen that the δ/ZrO2 interface is parallel with their habit planes,
ð111Þδ and (100)OII. In addition, some steps exist at the δ/ZrO2 interface. Dislocations are observed around these steps on the δ deuteride side. The formation of these steps and dislocations may be caused by the misfit between ð111Þδ and (100)OII because of their different interplanar spacings. From the ICDD crystallography data and the as-determined orientation relationship, a model of the transformation from δ deuteride to OIIZrO2 phase was proposed, as shown in Fig. 4. It shows that the ð111Þδ plane is parallel to the (100)OII plane. It can be seen that the ð111Þδ
Fig. 4. Model illustrating the transformation depending on the orientation relationship of ½011δ ==½011ZrO2 and ð111Þδ ==ð100ÞZrO2 . The green circles represent Zr atoms, gray circles represent D atoms, and red circles represent O atoms. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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C. Zhang et al. / Scripta Materialia 143 (2018) 1–4
Table 2 Calculated volume expansion and lattice misfit in different ZrO2 phases. Phase
Structure
δ-ZrH1.66 Fm3m P21/c P42/nmc Fm3m Pbcm Pnam
m-ZrO2 t-ZrO2 c-ZrO2 OI-ZrO2 OII-ZrO2
V1/nm3
V2/nm3
Ω/%
Matched planesa
Interface misfit
PDF#
/
0.02725
/
/
/
65–6972
0.03650 0.03335 0.03371 0.03472 0.03022
/ / / / /
33.9 22.4 23.7 27.4 10.9
(111)δ//(111)m (111)δ//(011)t (111)δ//(111)c (111)δ//(002)OI (111)δ//(200)OII
3.9% 6.7% 6.9% 5.1% 2.1%
65–2357 50–1089 49–1642 37–1413 49–1746
a Matched planes of these zirconia are the planes with the closest interplanar distance with (111)δ planes according to the standard PDF cards.
plane could transform into the (100)OII plane by increasing the atomic distance along the [011]δ direction from 0.337 to 0.365 nm and decreasing the atomic distance along the [101]δ direction from 0.337 to 0.334 nm and simultaneously enlarging the angle between [011]δ and [101]δ from 60° to 62.8°. Furthermore, the diffusion of deuterium and oxygen is also required. The δ deuteride can then be transformed into the OII-ZrO2 phase. There are several kinds of zirconia. Under atmospheric pressure, stable zirconia is monoclinic m-ZrO2 (P21/c) up to 1440 K, then tetragonal t-ZrO2 (P42/nmc) from 1440 to 2950 K, and finally cubic c-ZrO2 (Fm3m) as the temperature becomes higher than 2950 K [17]. Two orthorhombic ZrO2, denoted as OI and OII, respectively, have been reported to form under higher pressure conditions [18,19]. OI-ZrO2 with a space group of Pbcm has been observed at pressures from 3.5 to 13 GPa; however, it can also be found in doped zirconia, films, and nanoparticles at atmospheric pressure [20–22]. OII structure ZrO2 (space group Pnam) is reported to be stable only at pressures higher than 12.5 GPa [17,18]. However, in this study, OII-ZrO2 phase was found in deuterated samples at atmospheric pressure. The formed OII-ZrO2 phase is not likely to be the surface artifact introduced by foil preparation because it is an acicular precipitate rather than an oxide film. From the observed OII-ZrO2 morphology and the orientation relationship between OII-ZrO2 and δ deuteride, as shown in Fig. 3, it is reasonable to consider that the formation of OII-ZrO2 phase depends on the formation of δ deuteride. During deuterium absorption, dissolved oxygen is expelled and aggregates on the interface of δ deuterides because of the growth of δ deuterides, which have a lower solubility of oxygen than zirconium matrix [23]. This provides a suitable oxygen atom concentration for the formation of zirconia. Furthermore, the formation of OII-ZrO2 may also be related to the volume strain energy. To evaluate the volume strain energy during the formation of zirconia, the atom volume strain Ω was calculated by the following equation [24]: Ω = (V1 − V2) / V2 ∗ 100%, where V1 and V2 are the atom volumes of Zr in ZrO2 and δ deuteride cells, respectively. The calculated atom volume V1 and the atom volume strain Ω of different ZrO2 structures are shown in Table 2. This demonstrates that OIIZrO2 has the lowest volume strain among all kinds of ZrO2 phases—10.9% as compared to 33.9%, 22.4%, 23.7%, and 27.4% for the m-ZrO2, t-ZrO2, c-ZrO2, and OI-ZrO2, respectively. The lowest Ω of OIIZrO2 represents the lowest volume strain energy that is preferential to precipitate zirconia from the δ-deuteride phase.
Furthermore, the misfit between the ð111Þδ and the (200)OII habit planes is small. From Table 2, it can be seen clearly that this is the lowest degree of misfit compared with the misfits between other ZrO2 phases and δ deuterides. Furthermore, by a slight readjustment of Zr atoms within the ð111Þδ habit plane, i.e., an increase in the atomic distance along the [011]δ direction and a decrease in the atomic distance along the [011]δ direction, accompanied by a tiny lattice rotation, as shown in Fig. 4, the FCC lattice of δ deuterides can be transformed into the Pnam lattice of OII-ZrO2 phase, as shown by our experimental results. In summary, the acicular cotunnite-type OII-ZrO2 was observed in Zr-4 alloy after deuterium absorption at 1173 K. It precipitated from the interface of δ deuterides, followed by the orientation relationship of ½011OII ==½011δ and ð100ÞOII ==ð111Þδ . Some steps existed at the δ/ ZrO2 interface, together with the appearance of dislocations in the δ deuteride side. The effects of oxygen content, volume strain, and crystal structure play critical roles in the formation of OII-ZrO2. The discovery of OII-ZrO2 phase after the deuterium absorption will be helpful in the understanding of the oxidation and corrosion behavior of zirconium alloys. This work was supported by the National Natural Science Foundation of China (Nos. 21171018 and 51271021), the Beijing Natural Science Foundation (No. 2162025), and the State Key Laboratory for Advanced Metals and Materials (No. 2016-T02). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]
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