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Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing
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Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing C. García-Rosales a,∗ , P. López-Ruiz a , S. Alvarez-Martín a , A. Calvo a , N. Ordás a , F. Koch b , J. Brinkmann b a b
CEIT and Tecnun (University of Navarra), E-20018 San Sebastian, Spain Max-Planck-Institut für Plasmaphysik (IPP), EURATOM Association, D-85748 Garching, Germany
a r t i c l e
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Article history: Available online xxx Keywords: Tungsten alloys Oxidation resistance Armour material Mechanical alloying HIP
a b s t r a c t Self-passivating tungsten based alloys are expected to provide a major safety advantage compared to pure tungsten when used as first wall armour of future fusion reactors, due to the formation of a protective oxide scale, preventing the formation of volatile and radioactive WO3 in case of a loss of coolant accident with simultaneous air ingress. In this work results of isothermal oxidations tests at 800 and 1000 ◦ C on bulk alloy WCr12Ti2.5 performed by thermogravimetric analysis (TGA) and by exposure to flowing air in a furnace are presented. In both cases a thin, dense Cr2 O3 layer is found at the outer surface, below which a Cr2 WO6 scale and Ti2 CrO5 layers alternating with WO3 are formed. The Cr2 O3 , Cr2 WO6 and Ti2 CrO5 scales act as protective barriers against fast inward O2− diffusion. The oxidation kinetics seems to be linear for the furnace exposure tests while for the TGA tests at 800 ◦ C the kinetics is first parabolic, transforming into linear after an initial phase. The linear oxidation rates are 2–3 orders of magnitude lower than for pure W. © 2014 Elsevier B.V. All rights reserved.
1. Introduction Tungsten is the reference material for the first wall (FW) armour of future fusion reactors such as DEMO. However, the use of tungsten represents an important safety concern in case of a lossof-coolant accident (LOCA) with simultaneous air ingress into the reactor vessel. In such a situation, temperatures up to 1200 ◦ C – depending on the power plant conceptual design – can be achieved in the in-vessel components within several tens of days due to the decay heat [1], which would lead to fast W oxidation with the release of volatile radioactively activated tungsten oxides [2]. The specific activity of W a few days after shutdown, following irradiation during 5 years in the first wall neutron spectrum of Model B PPCS, amounts to about 5 × 1015 Bq/kg [3]. Assuming an approximate area of the DEMO FW of 1000 m2 and a thickness of only 1 mm, and taking into account the density of W of 19.3 g/cm3 , the resulting activity in the W first wall would be of the order of 1020 Bq. Using the linear oxidation rate of W at 1000 ◦ C of 1.4 × 10−2 mg cm−2 s−1 [4] and taking into account the high volatility of WO3 , a release of the whole W of the FW armour would take place as a result of the
∗ Corresponding author. Tel.: +34 943 21 28 00; fax: +34 943 21 30 76. E-mail address: [email protected] (C. García-Rosales).
accident, if no self-shielding effects are considered. These data indicate the seriousness of the problem with W oxidation in case of a LOCA plus air ingress, and the relevance of the development of oxidation resistant, self-passivating tungsten-based alloys for fusion FW application. A possible way for avoiding this important safety issue is the addition to W of oxide forming alloying elements which should lead to the growth of a self-passivating layer at high temperature in presence of oxygen, preventing further W oxidation. During normal operation, the surface of this material will consist of W, owing to preferential sputtering of the alloying elements. In previous work [4–6], different W alloys have been manufactured at IPP via magnetron sputtering, demonstrating that thin films of composition WCr14Ti2 (in wt.%) have an oxidation rate about four orders of magnitude lower than pure W in the temperature range 600–1000 ◦ C, and exhibit excellent self-passivating behaviour due to the formation of a stable Cr2 O3 protective scale. Since thin films are not applicable for the FW armour because thicknesses of several mm are required for safe operation, W-Cr-Ti bulk alloys were manufactured at CEIT by powder metallurgy (mechanical alloying (MA) + Hot Isostatic Pressing (HIP)) [7,8], obtaining nearly 100% dense materials. In this paper, results of isothermal oxidations tests carried out on bulk alloy WCr12Ti2.5 are presented. The tests were performed
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Please cite this article in press as: C. García-Rosales, et al., Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing, Fusion Eng. Des. (2014), http://dx.doi.org/10.1016/j.fusengdes.2014.04.057
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both by thermogravimetric analysis (TGA) and by exposure to flowing air in a furnace; the results of both methods are compared and discussed in terms of oxidation rate as well as microstructure and composition of the formed oxide scale. 2. Experimental details The bulk alloy WCr12Ti2.5 was produced by 40 h MA of the elemental powders in a planetary ball mill followed by compaction, glass encapsulation and HIPing at 1300 ◦ C and 150 MPa for 1 h, after which a fully dense material was obtained. The oxygen and carbon content after HIPing was <0.2 and 0.073 wt.%, respectively. Further details of the manufacturing and properties of this alloy can be found in [8]. The isothermal oxidation tests were performed both discontinuously by exposure in a furnace under flowing air and continuously by TGA. The exposure tests were carried out at 800 and 1000 ◦ C in a horizontal furnace (Carbolite RHF 14/32) using a tube of grade 310 stainless steel (SS) for introducing different atmospheres during each cycle. Six and four samples were used for the tests at 800 and 1000 ◦ C, respectively, corresponding to different exposure times: 6, 12, 24, 35, 42 and 48 h for exposure at 800 ◦ C and 0.5, 1, 3 and 5 h for 1000 ◦ C. The sample size was ∼14 mm× 7 mm × 5 mm for the tests at 800 ◦ C and ∼7 mm × 7 mm × 3 mm for those at 1000 ◦ C. The samples were grinded and polish down to 6 m. All edges were chamfered to avoid preferential oxidation. The samples were placed inside the SS tube on a porous alumina cradle in such a way that the contact area between sample and cradle was as small as possible. While the furnace was brought to the test temperature, the tube outside the furnace was filled with a flowing atmosphere of pure Ar. Once the test temperature was achieved, the tube containing the sample was introduced in the furnace. When the sample temperature equaled the furnace temperature, the atmosphere was changed to synthetic air (80% N2 , 20% O2 ) with a flow rate of 3.3 l/min. The test time started when the Ar valve was closed and the one for air was opened. Once the envisaged test time was achieved, the atmosphere was changed from air to Ar to avoid further oxidation during cooling. At that moment the tube was taken out from the furnace and free cooled. The heating rate was estimated to be of the order of 100 ◦ C/min. The sample was weighted before and after the tests using a balance with a sensitivity of ±10 g. Before each test the surface area of the chamfered samples was measured with an optical microscope. The TGA isothermal oxidation tests were performed only at 800 ◦ C in a commercial setup (NETZSCH STA 449 F1 Jupiter) equipped with a thermobalance with a sensitivity of ±25 ng, and using synthetic air of 80% Ar and 20% O2 at a pressure of 1 atm. The sample size was ∼7 mm × 7 mm × 4 mm, and they were prepared in the same way as described previously for isothermal furnace exposure tests. The procedure is described in detail in [6]. After isothermal oxidation, the crystallographic phases of the oxide scales were analyzed by grazing incidence X-ray diffraction (GIXRD) with CuK␣ radiation using grazing incidence angles of the order of 4◦ . The microstructure of the oxide scales was observed by field-emission gun scanning electron microscopy (FEG-SEM) and energy dispersive X-ray spectroscopy (EDX) mapping. For the FEGSEM examinations, the samples were Ni-platted to prevent scale loss during metallographic preparation, embedded and polished. 3. Results and discussion At high temperature the oxidation kinetics of numerous metals obey a parabolic law [9] 2
(m/A) = kp · t
(1)
Fig. 1. Mass gain per area, , and square of the mass gain per area, , as a function of the furnace exposure time at 800 ◦ C (top) and 1000 ◦ C (bottom), and corresponding linear and parabolic fits.
where m/A is the mass gain per unit area during time t and kp is the parabolic rate constant. Such a law corresponds to an oxidation rate limited by diffusion through the formed protective scale. In other cases the oxidation rate is constant, indicating a linear kinetics: m/A = kl · t
(2)
where kl is the linear rate constant. This kinetics may result if the formed scale is highly porous, so that the gas phase mass transfer takes place within the pores and the scale thickness has no bearing on the oxidation rate [10]. To determine whether for the WCr12Ti2.5 alloy the oxidation kinetic law is linear or parabolic, the mass gain per unit area and the square of the mass gain per area are plotted in Fig. 1 vs. the furnace exposure time at 800 and 1000 ◦ C. From the corresponding fitting curves the rate constants can be determined. Even though four to six points (corresponding to four to six exposure times) per temperature are rather scarce to draw definitive conclusions, some tendencies can be evidenced. At 800 ◦ C the linear fit provides a clearly better correlation coefficient R2 than the linear fit. It has to be noted that the time recording, as mentioned above, starts when the Ar valve is closed and the one for air opened. This means that during the time needed to completely replace the Ar by air, lower mass gains will be measured. However, this time is partly compensated with the one needed to replace the air by Ar at the end of the cycle time so that the total mass gain should not be affected. Nevertheless, this procedure will result in an overall shift of the points towards longer times, but do not affect the slope of the curve and thus, the oxidation rate. At 1000 ◦ C both the linear and the parabolic fit provide R2 values >0.9, which may suggest a mixed oxidation kinetics. In this case it seems more reasonable to assume a linear law as for 800 ◦ C, even though the linear fit yields a negative time for zero mass gain, which makes no sense but could be due to the low amount of fitting points.
Please cite this article in press as: C. García-Rosales, et al., Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing, Fusion Eng. Des. (2014), http://dx.doi.org/10.1016/j.fusengdes.2014.04.057
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Fig. 2. Mass gain per area during isothermal tests at 800 ◦ C measured by TGA up to 24 h and up to 48 h. The mass gain per area obtained by furnace exposure at 800 ◦ C (Fig. 1, top) is included for comparison. The arrows indicate the points at which the scale probably cracks.
In Fig. 2 the mass gain per unit area measured by continuous isothermal TGA at 800◦ is shown as a function of time. The previous values obtained by furnace exposure tests at the same temperature are included for comparison. A first TGA run was performed up to 24 h, resulting in good agreement with the furnace exposure tests values except the already mentioned “dead time” probably related with the time recording. Up to about 10–12 h the behaviour is parabolic. Repetition of the TGA run with a new sample up to 48 h results in a linear oxidation law for times above 12 h, in agreement with the furnace exposure test values at 42 and 48 h. The TGA runs show very good reproducibility. The TGA test up to 48 h shows a “wavy” form, suggesting that the oxide scale cracks (position of the arrows), after what a new parabolic growth occurs until the scale cracks again, resulting in an overall linear growth. In Fig. 3 the linear oxidation rates of bulk WCr12Ti2.5 from furnace exposure tests obtained by the previous parabolic fits are
3
Fig. 3. Arrhenius plot of linear oxidation rates for pure W measured by TGA [4], bulk WCr12Ti2.5 alloy measured by furnace exposure, and thin film WCr14Ti2 alloy measured by TGA [6].
presented in an Arrhenius plot and compared with the corresponding kl values of pure W [4] and a thin film of similar composition [6], both measured by TGA. It can be observed that the kl values of bulk WCr12Ti2.5 are slightly higher than those of thin films and 2–3 orders of magnitude lower than those of pure W for the two temperatures studied. The different microstructure of the magnetron sputter deposited thin films compared to the bulk alloy has most probably an influence on the diffusion of the oxide forming elements. Nevertheless, the results for the bulk alloy at the two temperatures studied show a remarkably good agreement with the values for the thin film of similar composition. The GIXRD patterns of the oxide scale formed on bulk WCr12Ti2.5 after furnace exposure tests at 800 and 1000 ◦ C are shown in Fig. 4 for different oxidation times, including the comparison to database files [11]. The patterns are quite complex with significant peaks overlap. Besides the main phase of the
Fig. 4. GIXRD patterns of WCr12Ti2.5 samples after isothermal furnace oxidation at 800 ◦ C (left) and 1000 ◦ C (right) for different times. The lines denote the peak positions for Cr2 O3 , Cr2 WO6 and WO3 from the Powder Diffraction File Database [11]. The arrows indicate some of the main peaks of Ti2 CrO5 (PDF 01-079-0302 [11]). The dots indicate the peak positions of the main bulk alloy phase: bcc W with Cr in solid solution [8].
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Fig. 5. FEG-SEM cross-sections of WCr12Ti2.5 samples after isothermal furnace oxidation at 800 ◦ C for different times.
alloy – a bcc W rich phase with Cr in solid solution [8] denoted here by Wss and indicated by dots – at least four oxides seems to be present both at 800 and 1000 ◦ C: Cr2 O3 , Cr2 WO6 , Ti2 CrO5 and WO3 . For both temperatures the intensity of the Wss phase decreases with oxidation time due to the formation of thicker scales. Both Cr2 O3 and Cr2 WO6 are present already at the lowest times for both temperatures and their intensity do not change significantly with oxidation time (except for 800 ◦ C and 12 h, where abnormally fewer oxides are present for unknown reasons). Monoclinic WO3 and Ti2 CrO5 are also present at the lowest times for both temperatures.
WO3 exhibits relatively high intensity at 800 ◦ C for 24 h exposure while for 48 h the intensity decreases, probably because it is found at a larger depth; at 1000 ◦ C there are no large variations with time. Ti2 CrO5 shows larger intensity at 800 ◦ C and 48 h than for 1000 ◦ C at all times. The high intensity peak visible at 800 ◦ C at ∼36.25◦ results mainly from the overlap of Cr2 WO6 and Ti2 CrO5 . Figs. 5 and 6 show the cross sections of the scales formed after furnace oxidation for different times at 800 and 1000 ◦ C, respectively. The thicknesses of all scales are in line with the registered mass gains (Fig. 1). The thickness of the scale formed after
Fig. 6. FEG-SEM cross-sections of WCr12Ti2.5 samples after isothermal furnace oxidation at 1000 ◦ C for different times.
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Fig. 7. EDX mappings of a cross sectional image of WCr12Ti2.5 after 5 h isothermal furnace oxidation at 1000 ◦ C.
Fig. 8. FEG- SEM cross-section of a larger region of a WCr12Ti2.5 sample after isothermal furnace oxidation at 1000 ◦ C for 3 h.
oxidation at 800 ◦ C for 48 h is of the order of 15–20 m. The scales formed after the TGA tests at 800 ◦ C (not included) have very similar appearance to those of Figs. 5 and 6. The microstructure of the unoxidized bulk alloy can be barely appreciated in Figs. 5 and 6. It consists of three homogeneously distributed phases: a bcc W-rich main phase with Cr in solid solution (light grey), a bcc Cr-rich phases with W in solution (dark grey), and a minority -Ti phase (black) [8]. In Fig. 7, the EDX elemental mappings of the scale formed after 5 h furnace oxidation at 1000 ◦ C is depicted. These mappings together with the GIXRD patterns allow the identification of the composition of the oxide scales of Figs. 5 and 6. The dark most outer layer observed on all cross sections of Figs. 5 and 6 consist of a dense Cr and O rich (and W free) scale, identified as Cr2 O3 in the GIXRD patterns. Below this layer a region containing mainly Cr, W and O is present, which corresponds most probably to the Cr2 WO6 observed in the GIXRD patterns. Below this region, horizontal layers rich in Ti, Cr and O (most probably Ti2 CrO5 ) alternate with W- and O-containing layers which are Cr-depleted and can be identified as WO3 . A larger WO3 region is in contact with the bulk alloy. These regions can be clearly appreciated in Fig. 8. In the O mapping of Fig. 7 a relatively large porosity can be appreciated for the scales below the Cr2 O3 top layer; one can also observe O accumulation in the Cr-rich phase of the bulk alloy in the region close to the interface, indicating that some O has diffused inside the alloy. In this same region the alloy is Ti depleted. The thin Cr2 O3 layer at the outer surface of all scales of Figs. 5 and 6 exhibits a nearly constant thickness with increasing oxidation time, being ∼0.1–0.2 m for 800 ◦ C and ∼0.5 m for 1000 ◦ C. At 800 ◦ C the Cr2 WO6 region below the Cr2 O3 layer is not continuous for 6 and 12 h exposure, which explains the relatively large Wss peaks in the corresponding GIXRD patterns. This is also
the case for the scales at 1000 ◦ C after 30 min and 1 h exposure. After 35, 42 and 48 h exposure at 800 ◦ C, the scale below the thin Cr2 O3 layer forms a continuous region. For 3 and 5 h exposure at 1000 ◦ C there is a continuous scale all over the bulk alloy. One observes in all cases an irregular growth of the outer scale, while the interface between the oxidized region and the bulk alloy is much smoother (see Fig. 8). It is likely that the thin Cr2 O3 layer cracks due to compressive stresses related to the large molar volume of the oxide. This results in faster oxidation and growth of other oxides like Cr2 WO6 until the new Cr2 O3 is restored. This successive cracking may be the reason for the “wavy” form of Fig. 2, as mentioned above, and is most probably responsible for the observed linear instead of parabolic oxidation rate. For the larger exposure times at both temperatures a band of cavities oriented mainly parallel to the alloy interface are present especially at the WO3 region. In Fig. 8, it can be realized that the WO3 has a columnar structure, in agreement with the observations of Telu at al. [12] for the W-Cr system and Cifuentes et al. [13] for pure W oxidation. They also observe the presence of porosity parallel to the alloy interface in WO3 , and in [13] a mechanism for its formation is proposed. 4. Conclusions From the observations presented in this paper the following conclusions can be drawn: – A thin, dense Cr2 O3 layer is formed from the very beginning at both investigated temperatures at the outer surface, indicating that Cr3+ is the cation with fastest outward diffusion. Once the layer has reached a certain thickness (∼0.1–0.2 m for 800 ◦ C, ∼0.5 m for 1000 ◦ C), it cracks successively, accelerating oxidation and resulting in high scale roughness.
Please cite this article in press as: C. García-Rosales, et al., Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing, Fusion Eng. Des. (2014), http://dx.doi.org/10.1016/j.fusengdes.2014.04.057
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– Below the thin Cr2 O3 layer, a Cr2 WO6 scale forms, below which a Ti2 CrO5 layer alternates with WO3 , which growth downwards into the alloy. The Cr2 O3 , Cr2 WO6 and Ti2 CrO5 scales act as protective barriers against fast inward O2− diffusion. – The oxidation kinetics seems to be linear at both temperatures for the furnace exposure tests while for the TGA tests at 800 ◦ C the kinetics is first parabolic, transforming into linear after an initial phase. Successive cracking and restoring of the dense Cr2 O3 layer is most probably responsible for the apparent linear kinetics. The linear oxidation rates are 2–3 orders of magnitude lower than for pure W. More work is needed to understand the involved reactions based on thermodynamic analysis. Acknowledgments This work was funded by the European Community within the framework of EFDA WP08-09, WP10, WP11 and WP12 under the contract of EURATOM/CIEMAT, as well as by the Spanish Ministry for Economy and Competitiveness (ENE2012-30753) and the Basque Government (ETORTEK-ACTIMAT 2013). Technical support of I. Andueza, CEIT, is greatly acknowledged.
References [1] D. Maisonnier, I. Cook, P. Sardain, R. Andreani, L. Di Pace, R. Forrest, et al., A Conceptual Study of Commercial Fusion Power Plants, Final Report, EFDA-RPRE-5.0, 13 April 2005. [2] N.P. Taylor, R. Pampin, Fusion Eng. Des. 81 (2006) 1333–1338. [3] R.A. Forrest, A. Tabasso, C. Danani, S. Jakhar, A.K. Shaw, Handbook of Activation Data Calculated Using EASY-2007, UKAEA FUS 552, March 2009. [4] F. Koch, H. Bolt, Phys. Scr. T128 (2007) 100–105. [5] F. Koch, S. Köppl, H. Bolt, J. Nucl. Mater. 386–388 (2009) 572–574. [6] F. Koch, J. Brinkmann, S. Lindig, T.P. Mishra, Ch. Linsmeier, Phys. Scr. T145 (2011) 014019 (5 pp). [7] P. López-Ruiz, N. Ordás, S. Lindig, F. Koch, I. Iturriza, C. García-Rosales, Phys. Scr. T145 (2011) 014018 (5 pp). [8] P. López-Ruiz, N. Ordás, I. Iturriza, M. Walter, E. Gaganidze, S. Lindig, et al., J. Nucl. Mater. 442 (2013) S219–S224. [9] P. Sarrazin, A. Galerie, J. Fouletier, Mechanisms of High Temperature Corrosion: A Kinetic Approach, Tras Tech Publications, Zurich, 2008. [10] D.J. Young, High Temperature Oxidation and Corrosion of Metals, Elsevier, Oxford, 2008. [11] International Centre for Diffraction Data (ICDD), Powder Diffraction File Database, 2013 www.icdd.com [12] S. Telu, A. Patra, M. Sankaranarayana, R. Mitra, S.K. Pabi, J. Refract. Met. Hard Mater. 36 (2012) 191–203. [13] S.C. Cifuentes, M.A. Monge, P. Pérez, Corros. Sci. 57 (2012) 114–121.
Please cite this article in press as: C. García-Rosales, et al., Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing, Fusion Eng. Des. (2014), http://dx.doi.org/10.1016/j.fusengdes.2014.04.057
ARTICLE IN PRESS
FUSION-7491; No. of Pages 6
Fusion Engineering and Design xxx (2014) xxx–xxx
Contents lists available at ScienceDirect
Fusion Engineering and Design journal homepage: www.elsevier.com/locate/fusengdes
Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing C. García-Rosales a,∗ , P. López-Ruiz a , S. Alvarez-Martín a , A. Calvo a , N. Ordás a , F. Koch b , J. Brinkmann b a b
CEIT and Tecnun (University of Navarra), E-20018 San Sebastian, Spain Max-Planck-Institut für Plasmaphysik (IPP), EURATOM Association, D-85748 Garching, Germany
a r t i c l e
i n f o
Article history: Available online xxx Keywords: Tungsten alloys Oxidation resistance Armour material Mechanical alloying HIP
a b s t r a c t Self-passivating tungsten based alloys are expected to provide a major safety advantage compared to pure tungsten when used as first wall armour of future fusion reactors, due to the formation of a protective oxide scale, preventing the formation of volatile and radioactive WO3 in case of a loss of coolant accident with simultaneous air ingress. In this work results of isothermal oxidations tests at 800 and 1000 ◦ C on bulk alloy WCr12Ti2.5 performed by thermogravimetric analysis (TGA) and by exposure to flowing air in a furnace are presented. In both cases a thin, dense Cr2 O3 layer is found at the outer surface, below which a Cr2 WO6 scale and Ti2 CrO5 layers alternating with WO3 are formed. The Cr2 O3 , Cr2 WO6 and Ti2 CrO5 scales act as protective barriers against fast inward O2− diffusion. The oxidation kinetics seems to be linear for the furnace exposure tests while for the TGA tests at 800 ◦ C the kinetics is first parabolic, transforming into linear after an initial phase. The linear oxidation rates are 2–3 orders of magnitude lower than for pure W. © 2014 Elsevier B.V. All rights reserved.
1. Introduction Tungsten is the reference material for the first wall (FW) armour of future fusion reactors such as DEMO. However, the use of tungsten represents an important safety concern in case of a lossof-coolant accident (LOCA) with simultaneous air ingress into the reactor vessel. In such a situation, temperatures up to 1200 ◦ C – depending on the power plant conceptual design – can be achieved in the in-vessel components within several tens of days due to the decay heat [1], which would lead to fast W oxidation with the release of volatile radioactively activated tungsten oxides [2]. The specific activity of W a few days after shutdown, following irradiation during 5 years in the first wall neutron spectrum of Model B PPCS, amounts to about 5 × 1015 Bq/kg [3]. Assuming an approximate area of the DEMO FW of 1000 m2 and a thickness of only 1 mm, and taking into account the density of W of 19.3 g/cm3 , the resulting activity in the W first wall would be of the order of 1020 Bq. Using the linear oxidation rate of W at 1000 ◦ C of 1.4 × 10−2 mg cm−2 s−1 [4] and taking into account the high volatility of WO3 , a release of the whole W of the FW armour would take place as a result of the
∗ Corresponding author. Tel.: +34 943 21 28 00; fax: +34 943 21 30 76. E-mail address: [email protected] (C. García-Rosales).
accident, if no self-shielding effects are considered. These data indicate the seriousness of the problem with W oxidation in case of a LOCA plus air ingress, and the relevance of the development of oxidation resistant, self-passivating tungsten-based alloys for fusion FW application. A possible way for avoiding this important safety issue is the addition to W of oxide forming alloying elements which should lead to the growth of a self-passivating layer at high temperature in presence of oxygen, preventing further W oxidation. During normal operation, the surface of this material will consist of W, owing to preferential sputtering of the alloying elements. In previous work [4–6], different W alloys have been manufactured at IPP via magnetron sputtering, demonstrating that thin films of composition WCr14Ti2 (in wt.%) have an oxidation rate about four orders of magnitude lower than pure W in the temperature range 600–1000 ◦ C, and exhibit excellent self-passivating behaviour due to the formation of a stable Cr2 O3 protective scale. Since thin films are not applicable for the FW armour because thicknesses of several mm are required for safe operation, W-Cr-Ti bulk alloys were manufactured at CEIT by powder metallurgy (mechanical alloying (MA) + Hot Isostatic Pressing (HIP)) [7,8], obtaining nearly 100% dense materials. In this paper, results of isothermal oxidations tests carried out on bulk alloy WCr12Ti2.5 are presented. The tests were performed
http://dx.doi.org/10.1016/j.fusengdes.2014.04.057 0920-3796/© 2014 Elsevier B.V. All rights reserved.
Please cite this article in press as: C. García-Rosales, et al., Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing, Fusion Eng. Des. (2014), http://dx.doi.org/10.1016/j.fusengdes.2014.04.057
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both by thermogravimetric analysis (TGA) and by exposure to flowing air in a furnace; the results of both methods are compared and discussed in terms of oxidation rate as well as microstructure and composition of the formed oxide scale. 2. Experimental details The bulk alloy WCr12Ti2.5 was produced by 40 h MA of the elemental powders in a planetary ball mill followed by compaction, glass encapsulation and HIPing at 1300 ◦ C and 150 MPa for 1 h, after which a fully dense material was obtained. The oxygen and carbon content after HIPing was <0.2 and 0.073 wt.%, respectively. Further details of the manufacturing and properties of this alloy can be found in [8]. The isothermal oxidation tests were performed both discontinuously by exposure in a furnace under flowing air and continuously by TGA. The exposure tests were carried out at 800 and 1000 ◦ C in a horizontal furnace (Carbolite RHF 14/32) using a tube of grade 310 stainless steel (SS) for introducing different atmospheres during each cycle. Six and four samples were used for the tests at 800 and 1000 ◦ C, respectively, corresponding to different exposure times: 6, 12, 24, 35, 42 and 48 h for exposure at 800 ◦ C and 0.5, 1, 3 and 5 h for 1000 ◦ C. The sample size was ∼14 mm× 7 mm × 5 mm for the tests at 800 ◦ C and ∼7 mm × 7 mm × 3 mm for those at 1000 ◦ C. The samples were grinded and polish down to 6 m. All edges were chamfered to avoid preferential oxidation. The samples were placed inside the SS tube on a porous alumina cradle in such a way that the contact area between sample and cradle was as small as possible. While the furnace was brought to the test temperature, the tube outside the furnace was filled with a flowing atmosphere of pure Ar. Once the test temperature was achieved, the tube containing the sample was introduced in the furnace. When the sample temperature equaled the furnace temperature, the atmosphere was changed to synthetic air (80% N2 , 20% O2 ) with a flow rate of 3.3 l/min. The test time started when the Ar valve was closed and the one for air was opened. Once the envisaged test time was achieved, the atmosphere was changed from air to Ar to avoid further oxidation during cooling. At that moment the tube was taken out from the furnace and free cooled. The heating rate was estimated to be of the order of 100 ◦ C/min. The sample was weighted before and after the tests using a balance with a sensitivity of ±10 g. Before each test the surface area of the chamfered samples was measured with an optical microscope. The TGA isothermal oxidation tests were performed only at 800 ◦ C in a commercial setup (NETZSCH STA 449 F1 Jupiter) equipped with a thermobalance with a sensitivity of ±25 ng, and using synthetic air of 80% Ar and 20% O2 at a pressure of 1 atm. The sample size was ∼7 mm × 7 mm × 4 mm, and they were prepared in the same way as described previously for isothermal furnace exposure tests. The procedure is described in detail in [6]. After isothermal oxidation, the crystallographic phases of the oxide scales were analyzed by grazing incidence X-ray diffraction (GIXRD) with CuK␣ radiation using grazing incidence angles of the order of 4◦ . The microstructure of the oxide scales was observed by field-emission gun scanning electron microscopy (FEG-SEM) and energy dispersive X-ray spectroscopy (EDX) mapping. For the FEGSEM examinations, the samples were Ni-platted to prevent scale loss during metallographic preparation, embedded and polished. 3. Results and discussion At high temperature the oxidation kinetics of numerous metals obey a parabolic law [9] 2
(m/A) = kp · t
(1)
Fig. 1. Mass gain per area, , and square of the mass gain per area, , as a function of the furnace exposure time at 800 ◦ C (top) and 1000 ◦ C (bottom), and corresponding linear and parabolic fits.
where m/A is the mass gain per unit area during time t and kp is the parabolic rate constant. Such a law corresponds to an oxidation rate limited by diffusion through the formed protective scale. In other cases the oxidation rate is constant, indicating a linear kinetics: m/A = kl · t
(2)
where kl is the linear rate constant. This kinetics may result if the formed scale is highly porous, so that the gas phase mass transfer takes place within the pores and the scale thickness has no bearing on the oxidation rate [10]. To determine whether for the WCr12Ti2.5 alloy the oxidation kinetic law is linear or parabolic, the mass gain per unit area and the square of the mass gain per area are plotted in Fig. 1 vs. the furnace exposure time at 800 and 1000 ◦ C. From the corresponding fitting curves the rate constants can be determined. Even though four to six points (corresponding to four to six exposure times) per temperature are rather scarce to draw definitive conclusions, some tendencies can be evidenced. At 800 ◦ C the linear fit provides a clearly better correlation coefficient R2 than the linear fit. It has to be noted that the time recording, as mentioned above, starts when the Ar valve is closed and the one for air opened. This means that during the time needed to completely replace the Ar by air, lower mass gains will be measured. However, this time is partly compensated with the one needed to replace the air by Ar at the end of the cycle time so that the total mass gain should not be affected. Nevertheless, this procedure will result in an overall shift of the points towards longer times, but do not affect the slope of the curve and thus, the oxidation rate. At 1000 ◦ C both the linear and the parabolic fit provide R2 values >0.9, which may suggest a mixed oxidation kinetics. In this case it seems more reasonable to assume a linear law as for 800 ◦ C, even though the linear fit yields a negative time for zero mass gain, which makes no sense but could be due to the low amount of fitting points.
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Fig. 2. Mass gain per area during isothermal tests at 800 ◦ C measured by TGA up to 24 h and up to 48 h. The mass gain per area obtained by furnace exposure at 800 ◦ C (Fig. 1, top) is included for comparison. The arrows indicate the points at which the scale probably cracks.
In Fig. 2 the mass gain per unit area measured by continuous isothermal TGA at 800◦ is shown as a function of time. The previous values obtained by furnace exposure tests at the same temperature are included for comparison. A first TGA run was performed up to 24 h, resulting in good agreement with the furnace exposure tests values except the already mentioned “dead time” probably related with the time recording. Up to about 10–12 h the behaviour is parabolic. Repetition of the TGA run with a new sample up to 48 h results in a linear oxidation law for times above 12 h, in agreement with the furnace exposure test values at 42 and 48 h. The TGA runs show very good reproducibility. The TGA test up to 48 h shows a “wavy” form, suggesting that the oxide scale cracks (position of the arrows), after what a new parabolic growth occurs until the scale cracks again, resulting in an overall linear growth. In Fig. 3 the linear oxidation rates of bulk WCr12Ti2.5 from furnace exposure tests obtained by the previous parabolic fits are
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Fig. 3. Arrhenius plot of linear oxidation rates for pure W measured by TGA [4], bulk WCr12Ti2.5 alloy measured by furnace exposure, and thin film WCr14Ti2 alloy measured by TGA [6].
presented in an Arrhenius plot and compared with the corresponding kl values of pure W [4] and a thin film of similar composition [6], both measured by TGA. It can be observed that the kl values of bulk WCr12Ti2.5 are slightly higher than those of thin films and 2–3 orders of magnitude lower than those of pure W for the two temperatures studied. The different microstructure of the magnetron sputter deposited thin films compared to the bulk alloy has most probably an influence on the diffusion of the oxide forming elements. Nevertheless, the results for the bulk alloy at the two temperatures studied show a remarkably good agreement with the values for the thin film of similar composition. The GIXRD patterns of the oxide scale formed on bulk WCr12Ti2.5 after furnace exposure tests at 800 and 1000 ◦ C are shown in Fig. 4 for different oxidation times, including the comparison to database files [11]. The patterns are quite complex with significant peaks overlap. Besides the main phase of the
Fig. 4. GIXRD patterns of WCr12Ti2.5 samples after isothermal furnace oxidation at 800 ◦ C (left) and 1000 ◦ C (right) for different times. The lines denote the peak positions for Cr2 O3 , Cr2 WO6 and WO3 from the Powder Diffraction File Database [11]. The arrows indicate some of the main peaks of Ti2 CrO5 (PDF 01-079-0302 [11]). The dots indicate the peak positions of the main bulk alloy phase: bcc W with Cr in solid solution [8].
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Fig. 5. FEG-SEM cross-sections of WCr12Ti2.5 samples after isothermal furnace oxidation at 800 ◦ C for different times.
alloy – a bcc W rich phase with Cr in solid solution [8] denoted here by Wss and indicated by dots – at least four oxides seems to be present both at 800 and 1000 ◦ C: Cr2 O3 , Cr2 WO6 , Ti2 CrO5 and WO3 . For both temperatures the intensity of the Wss phase decreases with oxidation time due to the formation of thicker scales. Both Cr2 O3 and Cr2 WO6 are present already at the lowest times for both temperatures and their intensity do not change significantly with oxidation time (except for 800 ◦ C and 12 h, where abnormally fewer oxides are present for unknown reasons). Monoclinic WO3 and Ti2 CrO5 are also present at the lowest times for both temperatures.
WO3 exhibits relatively high intensity at 800 ◦ C for 24 h exposure while for 48 h the intensity decreases, probably because it is found at a larger depth; at 1000 ◦ C there are no large variations with time. Ti2 CrO5 shows larger intensity at 800 ◦ C and 48 h than for 1000 ◦ C at all times. The high intensity peak visible at 800 ◦ C at ∼36.25◦ results mainly from the overlap of Cr2 WO6 and Ti2 CrO5 . Figs. 5 and 6 show the cross sections of the scales formed after furnace oxidation for different times at 800 and 1000 ◦ C, respectively. The thicknesses of all scales are in line with the registered mass gains (Fig. 1). The thickness of the scale formed after
Fig. 6. FEG-SEM cross-sections of WCr12Ti2.5 samples after isothermal furnace oxidation at 1000 ◦ C for different times.
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Fig. 7. EDX mappings of a cross sectional image of WCr12Ti2.5 after 5 h isothermal furnace oxidation at 1000 ◦ C.
Fig. 8. FEG- SEM cross-section of a larger region of a WCr12Ti2.5 sample after isothermal furnace oxidation at 1000 ◦ C for 3 h.
oxidation at 800 ◦ C for 48 h is of the order of 15–20 m. The scales formed after the TGA tests at 800 ◦ C (not included) have very similar appearance to those of Figs. 5 and 6. The microstructure of the unoxidized bulk alloy can be barely appreciated in Figs. 5 and 6. It consists of three homogeneously distributed phases: a bcc W-rich main phase with Cr in solid solution (light grey), a bcc Cr-rich phases with W in solution (dark grey), and a minority -Ti phase (black) [8]. In Fig. 7, the EDX elemental mappings of the scale formed after 5 h furnace oxidation at 1000 ◦ C is depicted. These mappings together with the GIXRD patterns allow the identification of the composition of the oxide scales of Figs. 5 and 6. The dark most outer layer observed on all cross sections of Figs. 5 and 6 consist of a dense Cr and O rich (and W free) scale, identified as Cr2 O3 in the GIXRD patterns. Below this layer a region containing mainly Cr, W and O is present, which corresponds most probably to the Cr2 WO6 observed in the GIXRD patterns. Below this region, horizontal layers rich in Ti, Cr and O (most probably Ti2 CrO5 ) alternate with W- and O-containing layers which are Cr-depleted and can be identified as WO3 . A larger WO3 region is in contact with the bulk alloy. These regions can be clearly appreciated in Fig. 8. In the O mapping of Fig. 7 a relatively large porosity can be appreciated for the scales below the Cr2 O3 top layer; one can also observe O accumulation in the Cr-rich phase of the bulk alloy in the region close to the interface, indicating that some O has diffused inside the alloy. In this same region the alloy is Ti depleted. The thin Cr2 O3 layer at the outer surface of all scales of Figs. 5 and 6 exhibits a nearly constant thickness with increasing oxidation time, being ∼0.1–0.2 m for 800 ◦ C and ∼0.5 m for 1000 ◦ C. At 800 ◦ C the Cr2 WO6 region below the Cr2 O3 layer is not continuous for 6 and 12 h exposure, which explains the relatively large Wss peaks in the corresponding GIXRD patterns. This is also
the case for the scales at 1000 ◦ C after 30 min and 1 h exposure. After 35, 42 and 48 h exposure at 800 ◦ C, the scale below the thin Cr2 O3 layer forms a continuous region. For 3 and 5 h exposure at 1000 ◦ C there is a continuous scale all over the bulk alloy. One observes in all cases an irregular growth of the outer scale, while the interface between the oxidized region and the bulk alloy is much smoother (see Fig. 8). It is likely that the thin Cr2 O3 layer cracks due to compressive stresses related to the large molar volume of the oxide. This results in faster oxidation and growth of other oxides like Cr2 WO6 until the new Cr2 O3 is restored. This successive cracking may be the reason for the “wavy” form of Fig. 2, as mentioned above, and is most probably responsible for the observed linear instead of parabolic oxidation rate. For the larger exposure times at both temperatures a band of cavities oriented mainly parallel to the alloy interface are present especially at the WO3 region. In Fig. 8, it can be realized that the WO3 has a columnar structure, in agreement with the observations of Telu at al. [12] for the W-Cr system and Cifuentes et al. [13] for pure W oxidation. They also observe the presence of porosity parallel to the alloy interface in WO3 , and in [13] a mechanism for its formation is proposed. 4. Conclusions From the observations presented in this paper the following conclusions can be drawn: – A thin, dense Cr2 O3 layer is formed from the very beginning at both investigated temperatures at the outer surface, indicating that Cr3+ is the cation with fastest outward diffusion. Once the layer has reached a certain thickness (∼0.1–0.2 m for 800 ◦ C, ∼0.5 m for 1000 ◦ C), it cracks successively, accelerating oxidation and resulting in high scale roughness.
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– Below the thin Cr2 O3 layer, a Cr2 WO6 scale forms, below which a Ti2 CrO5 layer alternates with WO3 , which growth downwards into the alloy. The Cr2 O3 , Cr2 WO6 and Ti2 CrO5 scales act as protective barriers against fast inward O2− diffusion. – The oxidation kinetics seems to be linear at both temperatures for the furnace exposure tests while for the TGA tests at 800 ◦ C the kinetics is first parabolic, transforming into linear after an initial phase. Successive cracking and restoring of the dense Cr2 O3 layer is most probably responsible for the apparent linear kinetics. The linear oxidation rates are 2–3 orders of magnitude lower than for pure W. More work is needed to understand the involved reactions based on thermodynamic analysis. Acknowledgments This work was funded by the European Community within the framework of EFDA WP08-09, WP10, WP11 and WP12 under the contract of EURATOM/CIEMAT, as well as by the Spanish Ministry for Economy and Competitiveness (ENE2012-30753) and the Basque Government (ETORTEK-ACTIMAT 2013). Technical support of I. Andueza, CEIT, is greatly acknowledged.
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Please cite this article in press as: C. García-Rosales, et al., Oxidation behaviour of bulk W-Cr-Ti alloys prepared by mechanical alloying and HIPing, Fusion Eng. Des. (2014), http://dx.doi.org/10.1016/j.fusengdes.2014.04.057