- Email: [email protected]
Microstructure, oxidation resistance and tensile properties of silicide coated Nb-alloy C-103
Materials Science & Engineering A 645 (2015) 339–346
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure, oxidation resistance and tensile properties of silicide coated Nb-alloy C-103 M. Sankar n, V.V. Satya Prasad, R.G. Baligidad, Md. Z. Alam, D.K. Das, A.A. Gokhale Defence Metallurgical Research Laboratory, Hyderabad 500058, Telangana, India
art ic l e i nf o
a b s t r a c t
Article history: Received 14 March 2015 Received in revised form 20 July 2015 Accepted 21 July 2015 Available online 22 July 2015
A Fe-Cr-alloyed silicide coating has been formed on a commercially available Nb-alloy called C-103 by using a slurry-based technique followed by vacuum diffusion treatment at 1380 °C. The coating microstructure reveals a three-layer structure with the outer layer consisting of NbSi2 phase and the inner interdiffusion layer consisting of lower silicides such as Nb5Si3 and Nb3Si. The intermediate layer of the coating is comprised of a Fe–Cr alloyed Nb-silicide phase and NbSi2. The coating provides good shortterm protection to the substrate against high temperature oxidation in air at 1100 and 1300 °C. It is found that the presence of the coating increases the tensile strength of the alloy. & 2015 Elsevier B.V. All rights reserved.
Keywords: Niobium alloy Silicide coating Oxidation High temperature strength
1. Introduction The future aerospace applications such as advanced turbojet engines used in subsonic, supersonic and trans-atmospheric flights will require materials with ever increasing temperature and load bearing capabilities for improved performance under extreme environmental conditions. The state-of-the-art superalloys can be utilized only up to a maximum temperature of about 1150 °C. Therefore, there exists a need for viable materials that can be used at even higher temperatures (1150–2000 °C) [1]. Nb-base alloys provide attractive options in this regard in terms of melting temperatures, high temperature strength, room temperature ductility, low ductile to brittle transition temperature and good workability as compared to superalloys. However, their high temperature application has been severely limited by their poor oxidation resistance [2–5]. These alloys oxidize easily above about 250 °C with oxidation rate rising sharply above 500 °C [6]. Significant amount of niobium alloy development activities were undertaken during 1950–70, and several potential alloy compositions identified [7]. These alloys were extensively explored as candidate materials for nuclear applications and structural components in aircraft, space vehicles, and rockets [8,9]. Considerable work has been carried out on improving the high temperature oxidation resistance of Nballoys by modifying the alloy chemistry and also by using protective coatings [10,11]. Efforts towards enhancing oxidation resistance by adding alloying elements such as Al, Si and Ti has n
Corresponding author. E-mail address: [email protected] (M. Sankar).
http://dx.doi.org/10.1016/j.msea.2015.07.063 0921-5093/& 2015 Elsevier B.V. All rights reserved.
yielded limited success and often led to the deterioration of the mechanical properties of these alloys. Therefore, use of these alloys at high temperatures by applying suitable oxidation-resistant coatings is considered a viable option. Among the various coating materials examined, silicide coatings, particularly R512E with a nominal composition of Si–20Fe–20Cr (wt%), have been found most suitable for Nb-base alloys [12]. Further, mostly slurry-based methods have been adopted for applying these silicide coatings because these methods enable easy application of coating on large and geometrically complex parts [12,13]. Although silicide coating on niobium alloys have long been in use for high temperature applications, their microstructural details, oxidation properties and mechanical behavior are not well documented in the literature [13]. The present study examines the microstructural aspects of a Fe–Cr containing silicide coating deposited on a commercially available Nb-alloy C-103 through a slurry-based method. Oxidation protection provided by this coating at 1100 and 1300 °C in air under cyclic heating and cooling conditions has also been evaluated. The effect of coating on the tensile properties of the alloy has also been studied.
2. Experimental procedure Niobium-base alloy C-103 has been used as the substrate material in this study and its nominal composition is provided in Table. 1. This alloy was available in the form of cold rolled and annealed sheets (3 mm thick) from M/s TJTM, U.K. Tensile specimens having 6 mm gauge width and 50 mm gauge length were
340
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
Table 1 Chemical composition of the as-received C-103 alloy. Elements
Concentration (wt%)
C N O H Zr Hf Ti Nb
0.01 0.009 0.01 0.001 0.69 11 0.95 Bal
prepared from these sheets by electro-discharge wire cutting. Some of the tensile samples, along with a few 20 20 mm2 size coupons, were deposited with a silicide coating. Prior to coating deposition, the specimens were given a surface cleaning treatment by dipping them in a solution containing HNO3, HF and H2SO4 in water. Subsequently, they were rinsed in distilled water and acetone, and then dried in air. These samples were then deposited with a layer of green coating by dipping them in a slurry containing 60 wt% Si, 20 wt% Cr and 20 wt% Fe powders dispersed in an organic medium at room temperature (RT). The samples were withdrawn from the slurry and dried in air. These green coated samples were subsequently subjected to a vacuum diffusion treatment at 1380 °C for 2 h to develop the silicide coating. Cyclic oxidation of both uncoated and coated coupons was carried out in air at 1100 and 1300 °C. Each cycle in the cyclic oxidation test consisted of heating the samples at the desired temperature for 15 min followed by naturally cooling them outside the furnace at room temperature for 15 min. The cyclic oxidation time reported in this study represents the cumulative duration of exposure at the high temperature, i.e., 1100 or 1300 °C. Both coated and uncoated tensile samples were tested using Gleeble testing equipment at 600, 800, 900, 1100 and 1200 °C at an initial strain rate of 8 10 4 s 1. While the uncoated specimens were tested in a partial vacuum of 5 10 1 mbar maintained inside Gleeble, the coated specimens were tested in air. Phase identification in uncoated and coated specimens was carried out using X-ray diffraction (XRD). Cu Kα radiation was used for this purpose. Metallographically polished transverse sections of as-received and as-coated coupons were etched with a solution consisting of 30 ml HF, 15 ml HNO3, 15 ml HCl and 10 ml water for measuring grain size of the substrate in as-received and as-coated conditions by using a Leitz metalloplan optical microscope. Grain size was measured by using lineal intercept method [14]. Microstructural analysis of the as-received C-103 alloy was also carried out using a FEI G2 transmission electron microscope (TEM) attached with an energy dispersive X-ray spectroscope (EDS) operating at 200 kV. The TEM foils were prepared by ion milling in a GATAN'S precision polishing system operating at 4 kV. A Cameca SX-100 electron probe micro-analyzer (EPMA) operating at 20 kV was used for the elemental analysis in the coating. Coating crosssections and fracture surfaces of failed tensile samples were also observed by using a Leo 440i scanning electron microscope (SEM).
3. Results and discussion 3.1. Microstructure The C-103 substrate alloy in both as-received and as-coated conditions exhibits a single-phase microstructure consisting of equiaxed recrystallized grains (Fig. 1). The average grain size of the alloy in as-coated condition is 115 mm (Fig. 1b) which is coarser
Fig. 1. Optical micrograph of C-103 substrate alloy: (a) in as received condition and (b) after coating treatment.
than the grain size of 42 mm in as-received condition (Fig. 1a). This can be attributed to the grain coarsening due to high temperature diffusion heat treatment (at 1380 °C for 2 h) applied after green slurry deposition. The above diffusion heat treatment temperature is higher than the recrystallization temperature of C-103 alloy, i.e. 1200–1315 °C [15]. TEM bright field image of as-received C-103 alloy shows the presence of very fine (100 nm) precipitates (Fig. 2a) that were not revealed in optical microscopy. These precipitates were identified as HfO2 using selected area diffraction, as shown in Fig. 2b. The EDS analysis in TEM also confirmed them to be HfO2 by revealing the presence of Hf and oxygen in these precipitates (Fig. 2c). The crystal structure of these particles was determined to be cubic with lattice parameter of 4.6 Å, which was consistent with the reported structure of HfO2 [16]. Being highly reactive, some of the Hf present in the alloy reacts with the dissolved oxygen, leading to the formation of the observed HfO2 particles in the matrix.
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
341
Fig. 2. TEM analysis of HfO2 precipitates present in as received C-103 alloy: (a) bright field image, (b) SAD pattern from the particles of HfO2 along [011] zone axis, and (c) EDS analysis of precipitates showing presence of Hf and oxygen.
The cross-sectional microstructure of the coating is shown in Fig. 3. The coating is fairly uniform in thickness varying in the range 100–110 mm. The concentration profiles of all the major substrate and coating elements are presented in Fig. 4. The outer layer of the coating is primarily constituted of Nb and Si with their concentrations remaining constant at 32 and 68 at%, respectively, over the thickness of this layer. The atom ratio of Nb:Si being approximately 1:2 indicates that NbSi2 phase (JCPDS no. 72-1032, hexagonal) constitutes the outer layer. This phase also contains minor amounts of Hf (o2 at%), Cr (o 13 at%) and Fe (o 12 at%) in solid solution (Fig. 4). The presence of NbSi2 phase in the outer
layer is also confirmed from XRD analysis, as shown in Fig. 5. The concentrations of Nb and Si decrease and that of Fe, Cr and Hf increase significantly from the outer layer to the two intermediate layers. These intermediate layers, marked as L1 and L2 in Fig. 3a are approximately 30 and 18 mm, respectively, in thickness. The inner intermediate layer (L2) is constituted of a single silicide phase of approximate composition (in at%) 25Nb–18Fe–10Cr– 45Si–2Hf. From Fig. 3a, it is evident that the outer NbSi2 layer and L2 have extended into each other to form the two-phase layer L1. Therefore, it is expected that the layer L1 is constituted of NbSi2 and the above mentioned Fe–Cr containing silicide phase that
342
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
Fig. 5. X-ray diffractograms obtained from surface of uncoated and coated alloy.
Fig. 3. Cross-sectional microstructure of the coating in as coated condition (a) overall coating and (b) magnified view of L1, L2 and IDZ.
100
NbSi2
80
Concentration (at. %)
Nb
Coating
90
IDZ
L1 + L2
70 60 50
Si
Substrate
40 30 20
0
Fe
Cr
10
Hf 0
20
40
60
Ti 80
100
120
Depth ( µm) Fig. 4. EMPA concentration profiles of various elements across the coating.
constitutes L2. However, the Fe–Cr containing silicide phase in L2 has a composition range, as evident from the EPMA analysis (Fig. 4), which approximately varies from (in at%) 20Nb–14Fe–
19Cr–45Si–2Hf to 17Nb–12Fe–24Cr–45Si–2Hf. A columnar grain structure is revealed in L2, as typically shown in Fig. 3b. An increase in Nb content along with the corresponding sharp decreases in Fe and Cr contents is observed over the interdiffusion zone (IDZ), as seen from Fig. 4. The IDZ has a thickness of about 24 mm and consists of three sub-layers which can be seen in the magnified image presented in Fig. 3c. The outer sub-layer shows a finger-like structure consisting of alternate fingers of NbSi2 phase and the silicide phase constituting L2. Several Hf-rich bright particles are also present in this sub-layer. The second sub-layer has a Nb-rich matrix containing numerous Hf-containing precipitates. The inner (or the third) sub-layer is much richer in Nb than the intermediate sub-layer and contains fewer precipitates. From the EPMA composition data, the matrix phase for the second and third sub-layers is determined to be Nb5Si3 (approximate composition in at% 58Nb–39Si–3Hf) and Nb3Si (70Nb–26Si–3Hf), respectively. These phases contain negligible amounts (0.1–0.3 at%) of Fe and Cr. The thicknesses of the three sub-layers of IDZ are about 8, 4 and 13 mm, respectively. EPMA X-ray dot maps for Nb, Si, Fe, Cr and Hf taken over the entire coating thickness, as shown in Fig. 6a–f, also confirm the graded composition of as-formed coating. Some Hfrich precipitates are found near the surface of the coating (Fig. 6f) as well as throughout the outer layer. These are identified as HfSi2 (JCPDS no. 72-1201, orthorhombic) from XRD analysis. Several authors have reported the formation of silicide coating on Nb-base alloys by pack siliconizing process [17–21]. The typical microstructure of a pack silicide coating on Nb-alloy C-103 from in-house experiments is presented in Fig. 7 [21]. This coating consists of an outer NbSi2 layer, which constitutes the bulk of the coating, and an inner IDZ. The IDZ had two sub-layers and contain the lower silicide phase Nb5Si3 along with NbSi2. Although the microstructure of the present slurry deposited coating (Fig. 3) appears significantly different from that obtained by pack siliconizing method (Fig. 7), the general similarities between the two coatings can be easily identified. Firstly, NbSi2 phase constitutes the outer layer in both cases and the oxidation resistance of the coatings is derived primarily from this phase. However, the NbSi2 phase in the present slurry-based coating additionally contains some amounts of Fe and Cr in solid solution (Fig. 4). Secondly, the IDZ in both the cases has multiple sub-layers formed by the outward diffusion of Nb from the substrate during coating formation and contains lower silicide phases such as Nb5Si3 and Nb3Si along with NbSi2. The major difference between the two silicide coatings is the presence of a zone constituting of Fe–Cr containing Nb-silicide phases, i.e. L1 and L2, between the outer NbSi2 layer and the IDZ in case of slurry silicide coating (Fig. 4). The presence of Fe and Cr as the coating forming elements along with Si leads to the
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
343
Fig. 6. EMPA X-ray mapping of silicide coated alloy: (a) BSE image, (b) Cr map, (c) Fe map, (d) Hf map, (e) Nb map and (f) Si map.
Fig. 7. Microstructure of the pack silicide coating on C-103 alloy as reported by Alam et al. [21].
generation of such a zone and the formation of the alloy silicide L2 layer in the coating. Comparative phase diagram study reveals that there is a two-phase field, i.e. NbSi2 and (Nb, Fe, Cr)6Si3 (as found in L1) in the quarternary phase diagram of Nb–Fe–Cr–Si, while no such binary phase field exists adjacent to NbSi2 in binary Nb–Si phase diagram [22]. In terms of microstructure, the present coating is quite similar to that reported by Alam et al. [21], which was also formed by a similar slurry silicide method. The presence of a two-phase zone, as indicated by L1 in Fig. 3 was, however, not very apparent in the coating reported by Alam et al. [21]. The formation mechanism of the graded coating structure during the vacuum diffusion treatment, as indicated in Fig. 3 has already been reported [12,21]. It has been shown that the microstructural evolution in the coating during diffusion treatment in slurry silicide process is fundamentally different from that in case of pack silicide process. During pack siliciding, the NbSi2 layer of the coating grows in an inward manner, i.e. into the substrate starting from the initial substrate surface. Such inward movement of the growth front is sustained because of the predominant inward diffusion of the silicon atoms released from the pack at the sample surface through the NbSi2 layer. Such inward growth of the coating is made possible because of the gradual availability of Si atoms from the pack at the sample
surface during coating formation. In case of slurry silicide coating, however, all the silicon, along with Fe and Cr, is available on the sample surface in powder form prior to the beginning of coating formation. During vacuum diffusion treatment, these elements react among themselves as well as with the substrate elements to form the coating. As the duration of exposure at the diffusion temperature increases, continuous redistribution of various elements takes place leading to the development of the eventual microstructure after 2 h (Fig. 3) having a graded composition, as shown in Fig. 4. It has been shown that the IDZ of the coating develops very early during coating formation although its three sub-layer structure develops gradually over the entire diffusion duration of 2 h [12]. Several through-thickness cracks are also present in the coating, as evident in Figs. 3 and 7. These cracks are a typical feature of diffusion silicide coatings formed on Nb alloys and their generation can be ascribed to the difference in coefficients of thermal expansion (CTE) between the coating and the substrate. The NbSi2 phase constituting the coating has a higher CTE (7.3–11.7 mm °C 1) than that of the Nb substrate alloy (7–8 mm °C 1) [23]. Being extremely brittle, the silicide coating develops through-thickness cracks under the influence of the thermal mismatch stresses generated between the substrate and the coating as the coated substrate cools after the vacuum diffusion treatment at 1380 °C. Such cracks have also been reported by Alam et al. in case of pack silicide coatings formed on Nb, Mo and W [24,25]. Some amount of Kirkendall voids are also found throughout the coating (Fig. 3). These voids are generated because of unequal diffusive mass transfer between various coating elements (Fe, Cr and Si) and the substrate during coating formation, i.e. during vacuum diffusion treatment at 1380 °C. 3.2. Oxidation performance As mentioned earlier, both uncoated and coated specimens were exposed to cyclic oxidation exposure at 1100 and 1300 °C in air. Fig. 8 presents the weight change plot for these samples at both the temperatures. For comparison purpose, plots corresponding to samples with pack silicide coating, as reported by Alam et al. [21], have also been shown in the above figure. The bare alloy, as expected, has very poor oxidation resistance that is
344
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
Fig. 8. Weight change plot obtained for oxidation of uncoated and coated alloy at 1100 °C and 1300 °C.
Fig. 10. Variation in tensile properties of uncoated and coated alloy (a) strength (YS and UTS) and (b) ductility (% elongation).
Fig. 9. Coated sample after oxidation at 1300 °C for 2 h: (a) surface morphology showing the formation of a glassy silica layer and (b) the corresponding crosssection view.
reflected in the sharp weight loss observed very early during oxidation exposure (Fig. 8). The slurry silicide coated alloy exhibits much better oxidation performance at both 1100 and 1300 °C. For these coated samples, the exposure time before the net weight change becomes negative, i.e. the weight of the oxidized sample becomes lower than its initial weight prior to oxidation, is nearly 2 h at 1100 °C and somewhat higher at 4 h at 1300 °C. Thus, it is clearly evident that slurry silicide coating can offer good oxidation protection to C-103 alloy in air for short term high temperature cycling applications. From Fig. 8 it is apparent that the slurry silicide coating offers similar oxidation protection to C-103 alloy as pack silicide coating at 1300 °C, although the latter coating appears to survive much longer (up to about 6 h) than the former at
1100 °C. The protection in both the above coatings is derived from the formation of an oxidation resistant glassy silica (SiO2) layer on the coating surface upon exposure to high temperatures, as shown in Fig. 9a for the coated sample exposed at 1300 °C for 2 h. The corresponding cross-sectional view of the coating showing the silica layer is shown in Fig. 9b. Apart from covering the exposed surface, the silica film also seals the coating cracks, as evident from Fig. 9b. Such sealing of coating cracks, reduces the oxidation damage of the underlying substrate which otherwise would be caused by the ingress of oxidation through the cracks. Because of higher temperature, the coverage of the glassy silica layer over the exposed surface as well as sealing of the cracks by silica are much better at 1300 °C than at 1100 °C. Therefore, the coating provided better protection at 1300 °C than at 1100 °C, as evident from the weight change data shown in Fig. 9. These aspects have also been reported in several previous publications [21,26]. 3.3. Tensile properties The high temperature tensile properties of both silicide coated and uncoated alloy are shown in Fig. 10a and b. The 0.2% yield strength (YS) and ultimate tensile strength (UTS) of the alloy decrease with increase in temperature in both uncoated and coated conditions. The YS of the uncoated alloy gradually decreases from about 240 MPa at 600 °C to 150 MPa at 1200 °C. The corresponding decrease in coated alloy is minimal between 600 and 800 °C with the YS value remaining virtually constant at 300 MPa. Beyond 800 °C, the YS of the coated alloy gradually decreases with increase in test temperature up to 1200 °C, as evident from Fig. 10a.
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
345
Fig. 11. SEM fractograph showing fracture surface of specimens tested at (a) 600 °C and (b) 1200 °C.
Unlike YS, an initial rise in UTS of both uncoated and coated alloy is observed as the temperature increases from 600 to 800 °C. The UTS of the uncoated alloy rose from 348 MPa at 600 °C to about 370 MPa at 800 °C. The corresponding UTS values for the coated alloy are 410 and 430 MPa, respectively. Beyond 800 °C, the UTS decreases gradually with increase in test temperature, as evident from Fig. 10a. The tensile ductility of the alloy, measured in terms of % elongation, in both coated and uncoated conditions shows a slight initial drop (Fig. 10b) which corresponds to the increase in the UTS. Beyond 800 °C, the ductility gradually increases with its value at 1200 °C becoming 30% and 15% for the uncoated and coated alloy, respectively. Such increase in strength and corresponding decrease in ductility at about 800 °C has been reported for refractory metals and alloys [27]. This behavior has been attributed to dynamic strain aging of these materials during testing. It is evident from Fig. 10a and b that the tensile strength (UTS and YS) values of coated C-103 alloy are significantly higher than those of the uncoated alloy at all the test temperatures. For example, the YS values of the coated alloy at 600 °C and 1100 °C are 305 and 265 MPa, respectively. The corresponding YS values for the uncoated alloy are lower at 235 and 180 MPa, respectively. A similar trend is also evident in the UTS values (Fig. 10a). Fracture surfaces of all failed tensile specimens (both uncoated and coated) reveal essentially dimple features indicative of ductile mode of fracture (Fig. 11a and b). However, in case of samples tested in air, the fracture surface is found to be covered with a layer of oxide. The above mentioned increase in strength of C-103 alloy after the application of a silicide coating is contrary to the earlier reported trend that the presence of such coatings usually leads to a loss in alloy strength [28,29]. For example, Tavassoli who studied the effect of Si–20Cr–20Fe coating on mechanical properties of Nb– 10W–2Zr alloy and observed that coating decreases the strength of alloy [29]. Glushkpo et al. have also observed that the presence of MoSi2 coating on a Nb–W–Mo–Zr alloy (5V2MTs-2) caused a decrease in the strength of the alloy at all test temperatures ranging between RT and 1100 °C [30]. Although the increase in strength of the present alloy in presence of a slurry silicide coating is interesting, the reason for the same is not clear to the authors at the present time. Further studies are being planned to investigate this aspect in greater detail.
4. Conclusions Silicide coating on Nb-alloy C-103 alloy has been formed by using a slurry method. The coating consists of a continuous outer layer consisting of NbSi2 phase. The intermediate layer of the coating has two sub-layers with the outer sub-layer consisting of NbSi2 and a Fe-Cr-alloyed niobium silicide phase, and the inner
sub-layer consisting of only alloyed Nb-silicide phase. The inner layer is the interdiffusion zone consisting of lower silicide phases such as Nb5Si3 and Nb3Si. Significant coarsening of the substrate grain size occurs during the vacuum diffusion treatment associated with coating formation. The silicide coating is found to provide good short-term oxidation protection in air to the substrate alloy at 1100 and 1300 °C under cyclic conditions. Interestingly, the presence of the coating also leads to an enhancement of the tensile strength of the alloy at all test temperatures (600– 1100 °C).
Acknowledgments The authors are grateful to Defence Research and Development Organization, Ministry of Defence, New Delhi for the financial support in carrying out this research work. The authors extend their thanks to Dr. K. Sayta Prasad for carrying out TEM work. The authors also would like to thank officers and staff of ERG, HTCG and MBG for giving technical support to carry out this work.
References [1] C. Birla, M. Hoch, Metall. Trans. A 6 (1975) 1631–1634. [2] M. Vilasi, M. Francois, R. Podor, J. Steinmetz, J. Alloy. Compd. 264 (1998) 244–249. [3] Q. Zhao, Y. Yong-Si, Mater. Rev. 17 (2) (2003) 29–31. [4] J.R. Destefano, Int. J. Refract. Met. Hard Mater. 18 (5) (2000) 237–243. [5] R.O. Suzuki, M. Ishikawa, K. Ono, J. Alloy. Compd. 306 (2000) 280–285. [6] T.S. Sheasby, J. Electrochem. Soc. 1115 (7) (1968) 695–700. [7] J.R. Davis, P. Allen, ASM Metals Handbook—Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, 10, ASM International, Metals Park, Ohio, USA (1990), p. 2. [8] R.G. Frank, Refractory Metal Alloys, Plenum Press, New York, USA (1968), p. 325. [9] C.C. Vojcik, in: J.J. Stephens, I. Ahmad (Eds.), High Temperature Niobium Alloys, The Minerals and Material Society, Warrendale, Pennsylvania, USA, 1991, p. 1. [10] M. Fukumoto, Y. Matsumura, S. Hayashi, T. Narita, K. Sakamoto, A. Kasama, R. Tanaka, Mater. Trans. 44 (4) (2003) 731–735. [11] K. Matsumura, T. Koyanagi, T. Ohmi, M. Kudoh, Mater. Trans. 44 (5) (2003) 861–865. [12] M.D. Novak, C.G. Levi, Proceeding of IMECE 2007, Seattle, Washington, USA, Nov. 11–15 (2007). [13] S. Priceman, L. Sama, Electrochem. Technol. 6 (1968) 315–319. [14] ASTM. E112-10: Standard Test Method for Determining Average Grains Size, West Conshohocken (2010), p. 26. [15] M. Sankar, R.G. Baligidad, D.V.V. Satyanarayana, A.A. Gokhale, Mater. Sci. Eng. A 574 (2013) 104–112. [16] X. Zhao, D. Vanderbilt, Phys. Rev. B. 65 (2002) 1–4. [17] Y. Li, W. Soboyejo, R.A. Rapp, Metall. Mater. Trans. B 30 (3) (1999) 495–504. [18] R.W. Barlett, R.P. Gage, Trans. Metall. Soc. AIME 233 (1965) 968–978. [19] J. Guille, L. Matini, J. Mater. Sci. Lett. 7 (1988) 952–954. [20] B.V. Cockeram, Surf. Coat. Technol. 76–77 (1995) 20–27. [21] Md. Z. Alam, A.S. Rao, D.K. Das, Oxid. Met. 73 (2010) 513–530. [22] J.C. Zhao, M.R. Jackson, L.A. Peluso, Mater. Sci. Eng. A 372 (1–2) (2004) 21–27.
346
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
[23] S.J. Grisaffe, S.R. Levine, US Patent 3931447 (1976). [24] Md. Z. Alam, D.K. Das, J. Alloy. Compd. 487 (2009) 335–340. [25] S. Govindarajan, B. Mishra, D.L. Olson, J.J. Moore, J. Disam, Surf. Coat. Technol. 76–77 (1995) 7–13. [26] X. Larrong, Y.I. Darrqing, Y. Lei, C. Zhargang, Trans. Nonferr. Met. Soc. China 15 (1) (2005) 18–25.
[27] [28] [29] [30]
S.N. Nasser, W. Guo, Mater. Sci. Eng. A 284 (2000) 202–210. R. Perkins, E. Wright, J. Test. Eval. 2 (5) (1974) 7–15. A.A. Tavassoli, Weld. J. Res. Supply 52 (4) (1973) 168–172. P.I. Glushko, B.M. Shirokov, A.V. Shiyan, Powder Metall. Met. Ceram. 49 (9– 10) (2011) 616–618.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure, oxidation resistance and tensile properties of silicide coated Nb-alloy C-103 M. Sankar n, V.V. Satya Prasad, R.G. Baligidad, Md. Z. Alam, D.K. Das, A.A. Gokhale Defence Metallurgical Research Laboratory, Hyderabad 500058, Telangana, India
art ic l e i nf o
a b s t r a c t
Article history: Received 14 March 2015 Received in revised form 20 July 2015 Accepted 21 July 2015 Available online 22 July 2015
A Fe-Cr-alloyed silicide coating has been formed on a commercially available Nb-alloy called C-103 by using a slurry-based technique followed by vacuum diffusion treatment at 1380 °C. The coating microstructure reveals a three-layer structure with the outer layer consisting of NbSi2 phase and the inner interdiffusion layer consisting of lower silicides such as Nb5Si3 and Nb3Si. The intermediate layer of the coating is comprised of a Fe–Cr alloyed Nb-silicide phase and NbSi2. The coating provides good shortterm protection to the substrate against high temperature oxidation in air at 1100 and 1300 °C. It is found that the presence of the coating increases the tensile strength of the alloy. & 2015 Elsevier B.V. All rights reserved.
Keywords: Niobium alloy Silicide coating Oxidation High temperature strength
1. Introduction The future aerospace applications such as advanced turbojet engines used in subsonic, supersonic and trans-atmospheric flights will require materials with ever increasing temperature and load bearing capabilities for improved performance under extreme environmental conditions. The state-of-the-art superalloys can be utilized only up to a maximum temperature of about 1150 °C. Therefore, there exists a need for viable materials that can be used at even higher temperatures (1150–2000 °C) [1]. Nb-base alloys provide attractive options in this regard in terms of melting temperatures, high temperature strength, room temperature ductility, low ductile to brittle transition temperature and good workability as compared to superalloys. However, their high temperature application has been severely limited by their poor oxidation resistance [2–5]. These alloys oxidize easily above about 250 °C with oxidation rate rising sharply above 500 °C [6]. Significant amount of niobium alloy development activities were undertaken during 1950–70, and several potential alloy compositions identified [7]. These alloys were extensively explored as candidate materials for nuclear applications and structural components in aircraft, space vehicles, and rockets [8,9]. Considerable work has been carried out on improving the high temperature oxidation resistance of Nballoys by modifying the alloy chemistry and also by using protective coatings [10,11]. Efforts towards enhancing oxidation resistance by adding alloying elements such as Al, Si and Ti has n
Corresponding author. E-mail address: [email protected] (M. Sankar).
http://dx.doi.org/10.1016/j.msea.2015.07.063 0921-5093/& 2015 Elsevier B.V. All rights reserved.
yielded limited success and often led to the deterioration of the mechanical properties of these alloys. Therefore, use of these alloys at high temperatures by applying suitable oxidation-resistant coatings is considered a viable option. Among the various coating materials examined, silicide coatings, particularly R512E with a nominal composition of Si–20Fe–20Cr (wt%), have been found most suitable for Nb-base alloys [12]. Further, mostly slurry-based methods have been adopted for applying these silicide coatings because these methods enable easy application of coating on large and geometrically complex parts [12,13]. Although silicide coating on niobium alloys have long been in use for high temperature applications, their microstructural details, oxidation properties and mechanical behavior are not well documented in the literature [13]. The present study examines the microstructural aspects of a Fe–Cr containing silicide coating deposited on a commercially available Nb-alloy C-103 through a slurry-based method. Oxidation protection provided by this coating at 1100 and 1300 °C in air under cyclic heating and cooling conditions has also been evaluated. The effect of coating on the tensile properties of the alloy has also been studied.
2. Experimental procedure Niobium-base alloy C-103 has been used as the substrate material in this study and its nominal composition is provided in Table. 1. This alloy was available in the form of cold rolled and annealed sheets (3 mm thick) from M/s TJTM, U.K. Tensile specimens having 6 mm gauge width and 50 mm gauge length were
340
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
Table 1 Chemical composition of the as-received C-103 alloy. Elements
Concentration (wt%)
C N O H Zr Hf Ti Nb
0.01 0.009 0.01 0.001 0.69 11 0.95 Bal
prepared from these sheets by electro-discharge wire cutting. Some of the tensile samples, along with a few 20 20 mm2 size coupons, were deposited with a silicide coating. Prior to coating deposition, the specimens were given a surface cleaning treatment by dipping them in a solution containing HNO3, HF and H2SO4 in water. Subsequently, they were rinsed in distilled water and acetone, and then dried in air. These samples were then deposited with a layer of green coating by dipping them in a slurry containing 60 wt% Si, 20 wt% Cr and 20 wt% Fe powders dispersed in an organic medium at room temperature (RT). The samples were withdrawn from the slurry and dried in air. These green coated samples were subsequently subjected to a vacuum diffusion treatment at 1380 °C for 2 h to develop the silicide coating. Cyclic oxidation of both uncoated and coated coupons was carried out in air at 1100 and 1300 °C. Each cycle in the cyclic oxidation test consisted of heating the samples at the desired temperature for 15 min followed by naturally cooling them outside the furnace at room temperature for 15 min. The cyclic oxidation time reported in this study represents the cumulative duration of exposure at the high temperature, i.e., 1100 or 1300 °C. Both coated and uncoated tensile samples were tested using Gleeble testing equipment at 600, 800, 900, 1100 and 1200 °C at an initial strain rate of 8 10 4 s 1. While the uncoated specimens were tested in a partial vacuum of 5 10 1 mbar maintained inside Gleeble, the coated specimens were tested in air. Phase identification in uncoated and coated specimens was carried out using X-ray diffraction (XRD). Cu Kα radiation was used for this purpose. Metallographically polished transverse sections of as-received and as-coated coupons were etched with a solution consisting of 30 ml HF, 15 ml HNO3, 15 ml HCl and 10 ml water for measuring grain size of the substrate in as-received and as-coated conditions by using a Leitz metalloplan optical microscope. Grain size was measured by using lineal intercept method [14]. Microstructural analysis of the as-received C-103 alloy was also carried out using a FEI G2 transmission electron microscope (TEM) attached with an energy dispersive X-ray spectroscope (EDS) operating at 200 kV. The TEM foils were prepared by ion milling in a GATAN'S precision polishing system operating at 4 kV. A Cameca SX-100 electron probe micro-analyzer (EPMA) operating at 20 kV was used for the elemental analysis in the coating. Coating crosssections and fracture surfaces of failed tensile samples were also observed by using a Leo 440i scanning electron microscope (SEM).
3. Results and discussion 3.1. Microstructure The C-103 substrate alloy in both as-received and as-coated conditions exhibits a single-phase microstructure consisting of equiaxed recrystallized grains (Fig. 1). The average grain size of the alloy in as-coated condition is 115 mm (Fig. 1b) which is coarser
Fig. 1. Optical micrograph of C-103 substrate alloy: (a) in as received condition and (b) after coating treatment.
than the grain size of 42 mm in as-received condition (Fig. 1a). This can be attributed to the grain coarsening due to high temperature diffusion heat treatment (at 1380 °C for 2 h) applied after green slurry deposition. The above diffusion heat treatment temperature is higher than the recrystallization temperature of C-103 alloy, i.e. 1200–1315 °C [15]. TEM bright field image of as-received C-103 alloy shows the presence of very fine (100 nm) precipitates (Fig. 2a) that were not revealed in optical microscopy. These precipitates were identified as HfO2 using selected area diffraction, as shown in Fig. 2b. The EDS analysis in TEM also confirmed them to be HfO2 by revealing the presence of Hf and oxygen in these precipitates (Fig. 2c). The crystal structure of these particles was determined to be cubic with lattice parameter of 4.6 Å, which was consistent with the reported structure of HfO2 [16]. Being highly reactive, some of the Hf present in the alloy reacts with the dissolved oxygen, leading to the formation of the observed HfO2 particles in the matrix.
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
341
Fig. 2. TEM analysis of HfO2 precipitates present in as received C-103 alloy: (a) bright field image, (b) SAD pattern from the particles of HfO2 along [011] zone axis, and (c) EDS analysis of precipitates showing presence of Hf and oxygen.
The cross-sectional microstructure of the coating is shown in Fig. 3. The coating is fairly uniform in thickness varying in the range 100–110 mm. The concentration profiles of all the major substrate and coating elements are presented in Fig. 4. The outer layer of the coating is primarily constituted of Nb and Si with their concentrations remaining constant at 32 and 68 at%, respectively, over the thickness of this layer. The atom ratio of Nb:Si being approximately 1:2 indicates that NbSi2 phase (JCPDS no. 72-1032, hexagonal) constitutes the outer layer. This phase also contains minor amounts of Hf (o2 at%), Cr (o 13 at%) and Fe (o 12 at%) in solid solution (Fig. 4). The presence of NbSi2 phase in the outer
layer is also confirmed from XRD analysis, as shown in Fig. 5. The concentrations of Nb and Si decrease and that of Fe, Cr and Hf increase significantly from the outer layer to the two intermediate layers. These intermediate layers, marked as L1 and L2 in Fig. 3a are approximately 30 and 18 mm, respectively, in thickness. The inner intermediate layer (L2) is constituted of a single silicide phase of approximate composition (in at%) 25Nb–18Fe–10Cr– 45Si–2Hf. From Fig. 3a, it is evident that the outer NbSi2 layer and L2 have extended into each other to form the two-phase layer L1. Therefore, it is expected that the layer L1 is constituted of NbSi2 and the above mentioned Fe–Cr containing silicide phase that
342
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
Fig. 5. X-ray diffractograms obtained from surface of uncoated and coated alloy.
Fig. 3. Cross-sectional microstructure of the coating in as coated condition (a) overall coating and (b) magnified view of L1, L2 and IDZ.
100
NbSi2
80
Concentration (at. %)
Nb
Coating
90
IDZ
L1 + L2
70 60 50
Si
Substrate
40 30 20
0
Fe
Cr
10
Hf 0
20
40
60
Ti 80
100
120
Depth ( µm) Fig. 4. EMPA concentration profiles of various elements across the coating.
constitutes L2. However, the Fe–Cr containing silicide phase in L2 has a composition range, as evident from the EPMA analysis (Fig. 4), which approximately varies from (in at%) 20Nb–14Fe–
19Cr–45Si–2Hf to 17Nb–12Fe–24Cr–45Si–2Hf. A columnar grain structure is revealed in L2, as typically shown in Fig. 3b. An increase in Nb content along with the corresponding sharp decreases in Fe and Cr contents is observed over the interdiffusion zone (IDZ), as seen from Fig. 4. The IDZ has a thickness of about 24 mm and consists of three sub-layers which can be seen in the magnified image presented in Fig. 3c. The outer sub-layer shows a finger-like structure consisting of alternate fingers of NbSi2 phase and the silicide phase constituting L2. Several Hf-rich bright particles are also present in this sub-layer. The second sub-layer has a Nb-rich matrix containing numerous Hf-containing precipitates. The inner (or the third) sub-layer is much richer in Nb than the intermediate sub-layer and contains fewer precipitates. From the EPMA composition data, the matrix phase for the second and third sub-layers is determined to be Nb5Si3 (approximate composition in at% 58Nb–39Si–3Hf) and Nb3Si (70Nb–26Si–3Hf), respectively. These phases contain negligible amounts (0.1–0.3 at%) of Fe and Cr. The thicknesses of the three sub-layers of IDZ are about 8, 4 and 13 mm, respectively. EPMA X-ray dot maps for Nb, Si, Fe, Cr and Hf taken over the entire coating thickness, as shown in Fig. 6a–f, also confirm the graded composition of as-formed coating. Some Hfrich precipitates are found near the surface of the coating (Fig. 6f) as well as throughout the outer layer. These are identified as HfSi2 (JCPDS no. 72-1201, orthorhombic) from XRD analysis. Several authors have reported the formation of silicide coating on Nb-base alloys by pack siliconizing process [17–21]. The typical microstructure of a pack silicide coating on Nb-alloy C-103 from in-house experiments is presented in Fig. 7 [21]. This coating consists of an outer NbSi2 layer, which constitutes the bulk of the coating, and an inner IDZ. The IDZ had two sub-layers and contain the lower silicide phase Nb5Si3 along with NbSi2. Although the microstructure of the present slurry deposited coating (Fig. 3) appears significantly different from that obtained by pack siliconizing method (Fig. 7), the general similarities between the two coatings can be easily identified. Firstly, NbSi2 phase constitutes the outer layer in both cases and the oxidation resistance of the coatings is derived primarily from this phase. However, the NbSi2 phase in the present slurry-based coating additionally contains some amounts of Fe and Cr in solid solution (Fig. 4). Secondly, the IDZ in both the cases has multiple sub-layers formed by the outward diffusion of Nb from the substrate during coating formation and contains lower silicide phases such as Nb5Si3 and Nb3Si along with NbSi2. The major difference between the two silicide coatings is the presence of a zone constituting of Fe–Cr containing Nb-silicide phases, i.e. L1 and L2, between the outer NbSi2 layer and the IDZ in case of slurry silicide coating (Fig. 4). The presence of Fe and Cr as the coating forming elements along with Si leads to the
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
343
Fig. 6. EMPA X-ray mapping of silicide coated alloy: (a) BSE image, (b) Cr map, (c) Fe map, (d) Hf map, (e) Nb map and (f) Si map.
Fig. 7. Microstructure of the pack silicide coating on C-103 alloy as reported by Alam et al. [21].
generation of such a zone and the formation of the alloy silicide L2 layer in the coating. Comparative phase diagram study reveals that there is a two-phase field, i.e. NbSi2 and (Nb, Fe, Cr)6Si3 (as found in L1) in the quarternary phase diagram of Nb–Fe–Cr–Si, while no such binary phase field exists adjacent to NbSi2 in binary Nb–Si phase diagram [22]. In terms of microstructure, the present coating is quite similar to that reported by Alam et al. [21], which was also formed by a similar slurry silicide method. The presence of a two-phase zone, as indicated by L1 in Fig. 3 was, however, not very apparent in the coating reported by Alam et al. [21]. The formation mechanism of the graded coating structure during the vacuum diffusion treatment, as indicated in Fig. 3 has already been reported [12,21]. It has been shown that the microstructural evolution in the coating during diffusion treatment in slurry silicide process is fundamentally different from that in case of pack silicide process. During pack siliciding, the NbSi2 layer of the coating grows in an inward manner, i.e. into the substrate starting from the initial substrate surface. Such inward movement of the growth front is sustained because of the predominant inward diffusion of the silicon atoms released from the pack at the sample surface through the NbSi2 layer. Such inward growth of the coating is made possible because of the gradual availability of Si atoms from the pack at the sample
surface during coating formation. In case of slurry silicide coating, however, all the silicon, along with Fe and Cr, is available on the sample surface in powder form prior to the beginning of coating formation. During vacuum diffusion treatment, these elements react among themselves as well as with the substrate elements to form the coating. As the duration of exposure at the diffusion temperature increases, continuous redistribution of various elements takes place leading to the development of the eventual microstructure after 2 h (Fig. 3) having a graded composition, as shown in Fig. 4. It has been shown that the IDZ of the coating develops very early during coating formation although its three sub-layer structure develops gradually over the entire diffusion duration of 2 h [12]. Several through-thickness cracks are also present in the coating, as evident in Figs. 3 and 7. These cracks are a typical feature of diffusion silicide coatings formed on Nb alloys and their generation can be ascribed to the difference in coefficients of thermal expansion (CTE) between the coating and the substrate. The NbSi2 phase constituting the coating has a higher CTE (7.3–11.7 mm °C 1) than that of the Nb substrate alloy (7–8 mm °C 1) [23]. Being extremely brittle, the silicide coating develops through-thickness cracks under the influence of the thermal mismatch stresses generated between the substrate and the coating as the coated substrate cools after the vacuum diffusion treatment at 1380 °C. Such cracks have also been reported by Alam et al. in case of pack silicide coatings formed on Nb, Mo and W [24,25]. Some amount of Kirkendall voids are also found throughout the coating (Fig. 3). These voids are generated because of unequal diffusive mass transfer between various coating elements (Fe, Cr and Si) and the substrate during coating formation, i.e. during vacuum diffusion treatment at 1380 °C. 3.2. Oxidation performance As mentioned earlier, both uncoated and coated specimens were exposed to cyclic oxidation exposure at 1100 and 1300 °C in air. Fig. 8 presents the weight change plot for these samples at both the temperatures. For comparison purpose, plots corresponding to samples with pack silicide coating, as reported by Alam et al. [21], have also been shown in the above figure. The bare alloy, as expected, has very poor oxidation resistance that is
344
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
Fig. 8. Weight change plot obtained for oxidation of uncoated and coated alloy at 1100 °C and 1300 °C.
Fig. 10. Variation in tensile properties of uncoated and coated alloy (a) strength (YS and UTS) and (b) ductility (% elongation).
Fig. 9. Coated sample after oxidation at 1300 °C for 2 h: (a) surface morphology showing the formation of a glassy silica layer and (b) the corresponding crosssection view.
reflected in the sharp weight loss observed very early during oxidation exposure (Fig. 8). The slurry silicide coated alloy exhibits much better oxidation performance at both 1100 and 1300 °C. For these coated samples, the exposure time before the net weight change becomes negative, i.e. the weight of the oxidized sample becomes lower than its initial weight prior to oxidation, is nearly 2 h at 1100 °C and somewhat higher at 4 h at 1300 °C. Thus, it is clearly evident that slurry silicide coating can offer good oxidation protection to C-103 alloy in air for short term high temperature cycling applications. From Fig. 8 it is apparent that the slurry silicide coating offers similar oxidation protection to C-103 alloy as pack silicide coating at 1300 °C, although the latter coating appears to survive much longer (up to about 6 h) than the former at
1100 °C. The protection in both the above coatings is derived from the formation of an oxidation resistant glassy silica (SiO2) layer on the coating surface upon exposure to high temperatures, as shown in Fig. 9a for the coated sample exposed at 1300 °C for 2 h. The corresponding cross-sectional view of the coating showing the silica layer is shown in Fig. 9b. Apart from covering the exposed surface, the silica film also seals the coating cracks, as evident from Fig. 9b. Such sealing of coating cracks, reduces the oxidation damage of the underlying substrate which otherwise would be caused by the ingress of oxidation through the cracks. Because of higher temperature, the coverage of the glassy silica layer over the exposed surface as well as sealing of the cracks by silica are much better at 1300 °C than at 1100 °C. Therefore, the coating provided better protection at 1300 °C than at 1100 °C, as evident from the weight change data shown in Fig. 9. These aspects have also been reported in several previous publications [21,26]. 3.3. Tensile properties The high temperature tensile properties of both silicide coated and uncoated alloy are shown in Fig. 10a and b. The 0.2% yield strength (YS) and ultimate tensile strength (UTS) of the alloy decrease with increase in temperature in both uncoated and coated conditions. The YS of the uncoated alloy gradually decreases from about 240 MPa at 600 °C to 150 MPa at 1200 °C. The corresponding decrease in coated alloy is minimal between 600 and 800 °C with the YS value remaining virtually constant at 300 MPa. Beyond 800 °C, the YS of the coated alloy gradually decreases with increase in test temperature up to 1200 °C, as evident from Fig. 10a.
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
345
Fig. 11. SEM fractograph showing fracture surface of specimens tested at (a) 600 °C and (b) 1200 °C.
Unlike YS, an initial rise in UTS of both uncoated and coated alloy is observed as the temperature increases from 600 to 800 °C. The UTS of the uncoated alloy rose from 348 MPa at 600 °C to about 370 MPa at 800 °C. The corresponding UTS values for the coated alloy are 410 and 430 MPa, respectively. Beyond 800 °C, the UTS decreases gradually with increase in test temperature, as evident from Fig. 10a. The tensile ductility of the alloy, measured in terms of % elongation, in both coated and uncoated conditions shows a slight initial drop (Fig. 10b) which corresponds to the increase in the UTS. Beyond 800 °C, the ductility gradually increases with its value at 1200 °C becoming 30% and 15% for the uncoated and coated alloy, respectively. Such increase in strength and corresponding decrease in ductility at about 800 °C has been reported for refractory metals and alloys [27]. This behavior has been attributed to dynamic strain aging of these materials during testing. It is evident from Fig. 10a and b that the tensile strength (UTS and YS) values of coated C-103 alloy are significantly higher than those of the uncoated alloy at all the test temperatures. For example, the YS values of the coated alloy at 600 °C and 1100 °C are 305 and 265 MPa, respectively. The corresponding YS values for the uncoated alloy are lower at 235 and 180 MPa, respectively. A similar trend is also evident in the UTS values (Fig. 10a). Fracture surfaces of all failed tensile specimens (both uncoated and coated) reveal essentially dimple features indicative of ductile mode of fracture (Fig. 11a and b). However, in case of samples tested in air, the fracture surface is found to be covered with a layer of oxide. The above mentioned increase in strength of C-103 alloy after the application of a silicide coating is contrary to the earlier reported trend that the presence of such coatings usually leads to a loss in alloy strength [28,29]. For example, Tavassoli who studied the effect of Si–20Cr–20Fe coating on mechanical properties of Nb– 10W–2Zr alloy and observed that coating decreases the strength of alloy [29]. Glushkpo et al. have also observed that the presence of MoSi2 coating on a Nb–W–Mo–Zr alloy (5V2MTs-2) caused a decrease in the strength of the alloy at all test temperatures ranging between RT and 1100 °C [30]. Although the increase in strength of the present alloy in presence of a slurry silicide coating is interesting, the reason for the same is not clear to the authors at the present time. Further studies are being planned to investigate this aspect in greater detail.
4. Conclusions Silicide coating on Nb-alloy C-103 alloy has been formed by using a slurry method. The coating consists of a continuous outer layer consisting of NbSi2 phase. The intermediate layer of the coating has two sub-layers with the outer sub-layer consisting of NbSi2 and a Fe-Cr-alloyed niobium silicide phase, and the inner
sub-layer consisting of only alloyed Nb-silicide phase. The inner layer is the interdiffusion zone consisting of lower silicide phases such as Nb5Si3 and Nb3Si. Significant coarsening of the substrate grain size occurs during the vacuum diffusion treatment associated with coating formation. The silicide coating is found to provide good short-term oxidation protection in air to the substrate alloy at 1100 and 1300 °C under cyclic conditions. Interestingly, the presence of the coating also leads to an enhancement of the tensile strength of the alloy at all test temperatures (600– 1100 °C).
Acknowledgments The authors are grateful to Defence Research and Development Organization, Ministry of Defence, New Delhi for the financial support in carrying out this research work. The authors extend their thanks to Dr. K. Sayta Prasad for carrying out TEM work. The authors also would like to thank officers and staff of ERG, HTCG and MBG for giving technical support to carry out this work.
References [1] C. Birla, M. Hoch, Metall. Trans. A 6 (1975) 1631–1634. [2] M. Vilasi, M. Francois, R. Podor, J. Steinmetz, J. Alloy. Compd. 264 (1998) 244–249. [3] Q. Zhao, Y. Yong-Si, Mater. Rev. 17 (2) (2003) 29–31. [4] J.R. Destefano, Int. J. Refract. Met. Hard Mater. 18 (5) (2000) 237–243. [5] R.O. Suzuki, M. Ishikawa, K. Ono, J. Alloy. Compd. 306 (2000) 280–285. [6] T.S. Sheasby, J. Electrochem. Soc. 1115 (7) (1968) 695–700. [7] J.R. Davis, P. Allen, ASM Metals Handbook—Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, 10, ASM International, Metals Park, Ohio, USA (1990), p. 2. [8] R.G. Frank, Refractory Metal Alloys, Plenum Press, New York, USA (1968), p. 325. [9] C.C. Vojcik, in: J.J. Stephens, I. Ahmad (Eds.), High Temperature Niobium Alloys, The Minerals and Material Society, Warrendale, Pennsylvania, USA, 1991, p. 1. [10] M. Fukumoto, Y. Matsumura, S. Hayashi, T. Narita, K. Sakamoto, A. Kasama, R. Tanaka, Mater. Trans. 44 (4) (2003) 731–735. [11] K. Matsumura, T. Koyanagi, T. Ohmi, M. Kudoh, Mater. Trans. 44 (5) (2003) 861–865. [12] M.D. Novak, C.G. Levi, Proceeding of IMECE 2007, Seattle, Washington, USA, Nov. 11–15 (2007). [13] S. Priceman, L. Sama, Electrochem. Technol. 6 (1968) 315–319. [14] ASTM. E112-10: Standard Test Method for Determining Average Grains Size, West Conshohocken (2010), p. 26. [15] M. Sankar, R.G. Baligidad, D.V.V. Satyanarayana, A.A. Gokhale, Mater. Sci. Eng. A 574 (2013) 104–112. [16] X. Zhao, D. Vanderbilt, Phys. Rev. B. 65 (2002) 1–4. [17] Y. Li, W. Soboyejo, R.A. Rapp, Metall. Mater. Trans. B 30 (3) (1999) 495–504. [18] R.W. Barlett, R.P. Gage, Trans. Metall. Soc. AIME 233 (1965) 968–978. [19] J. Guille, L. Matini, J. Mater. Sci. Lett. 7 (1988) 952–954. [20] B.V. Cockeram, Surf. Coat. Technol. 76–77 (1995) 20–27. [21] Md. Z. Alam, A.S. Rao, D.K. Das, Oxid. Met. 73 (2010) 513–530. [22] J.C. Zhao, M.R. Jackson, L.A. Peluso, Mater. Sci. Eng. A 372 (1–2) (2004) 21–27.
346
M. Sankar et al. / Materials Science & Engineering A 645 (2015) 339–346
[23] S.J. Grisaffe, S.R. Levine, US Patent 3931447 (1976). [24] Md. Z. Alam, D.K. Das, J. Alloy. Compd. 487 (2009) 335–340. [25] S. Govindarajan, B. Mishra, D.L. Olson, J.J. Moore, J. Disam, Surf. Coat. Technol. 76–77 (1995) 7–13. [26] X. Larrong, Y.I. Darrqing, Y. Lei, C. Zhargang, Trans. Nonferr. Met. Soc. China 15 (1) (2005) 18–25.
[27] [28] [29] [30]
S.N. Nasser, W. Guo, Mater. Sci. Eng. A 284 (2000) 202–210. R. Perkins, E. Wright, J. Test. Eval. 2 (5) (1974) 7–15. A.A. Tavassoli, Weld. J. Res. Supply 52 (4) (1973) 168–172. P.I. Glushko, B.M. Shirokov, A.V. Shiyan, Powder Metall. Met. Ceram. 49 (9– 10) (2011) 616–618.