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Quaternary chromium-based alloys strengthened by Heusler phase precipitation
Materials Science & Engineering A 647 (2015) 322–332
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Quaternary chromium-based alloys strengthened by Heusler phase precipitation D. Locq n, P. Caron, C. Ramusat, R. Mévrel Onera – The French Aerospace Lab, F-92322 Châtillon, France
art ic l e i nf o Article history: Received 29 July 2015 Received in revised form 7 September 2015 Accepted 8 September 2015 Available online 11 September 2015 Keywords: Chromium alloy Heusler-phase precipitation Electron microscopy Microanalysis Oxidation Creep
a b s t r a c t Five quaternary chromium-based alloys have been designed in order to precipitate β′-Ni2AlTi Heusler phase strengthening particles while forming a protective alumina scale. These alloys were prepared by arc or induction melting as small polycrystalline ingots. β′-Ni2AlTi Heusler phase was observed as massive particles in the interdendritic areas of the cast alloys and as fine precipitates within the dendrites. The level of aluminium added in these alloys was however not enough to promote the formation of a continuous alumina protective scale before or during cyclic oxidation tests at 1100 °C. Surface aluminium enrichment of cyclic oxidation samples was carried out by vapour-phase aluminization to promote formation of a protective alumina scale at 1100 °C. That resulted in a dramatic improvement of the cyclic oxidation resistance of the Cr-based alloys. Compressive creep test performed at 850 °C and 1100 °C demonstrated the significant strengthening effect of the β′-Ni2AlTi Heusler phase precipitates when compared with pure chromium. & 2015 Elsevier B.V. All rights reserved.
1. Introduction Nickel-based superalloys are used for decades as materials for high temperature aero-engine components such as blades, vanes, disks and casings. Development of superalloys with improved performance is becoming more and more difficult and costly and a new challenge for aero-engine manufacturers is to try to replace some of them with less dense and/or more mechanical and environmental resistant materials. This pressing goal is motivated by economical and performance requirements. Chromium-based alloys have been considered as potential candidates for high-temperature applications since the late 1940's because of their high melting point, high elastic modulus, low density and of the good oxidation resistance of these chromiaforming materials [1–3]. During the 1950's and 1960's, extensive research projects were mainly conducted in Australia, in the United States of America and in the Soviet Union [1]. However, these works highlighted or confirmed that chromium-based alloys could present some drawbacks, such as a high ductile-brittle transition temperature (DBTT) and a high nitrogen solubility which can cause severe embrittlement after exposure at high temperature in air. Moreover, it has not been proved till today that very good high-temperature mechanical properties could be reached. Given n Correspondence to: Onera, DMSM, BP72, 29 avenue de la Division Leclerc, 92322 Châtillon Cedex, France. Fax: þ33 146734164. E-mail address: [email protected] (D. Locq).
http://dx.doi.org/10.1016/j.msea.2015.09.033 0921-5093/& 2015 Elsevier B.V. All rights reserved.
these difficulties and the concomitant, rapid and successful development of nickel-based superalloys, effort to develop high temperature chromium-based alloys ceased in the late 1970's. However, research programs on such alloys have recently resumed. The potential applications are for general high temperature structural materials [3–8], or more specifically gas turbines [9–12], for nuclear plants [13,14] as well as for glass tools [15]. The present study aims at designing new chromium-based alloys with improved high-temperature mechanical and environmental resistance in order to propose an alternative to nickelbased superalloys for high temperature component applications in gas turbine engines, such as turbine vanes and blades. The alloys investigated here are strengthened by precipitates of the intermetallic β′-Ni2AlTi compound and are designed for developing a protective alumina scale instead of the chromia scale which usually forms on chromium-based alloys but which is unstable in the high temperature regime (T4 900 °C).
2. Alloy design Components in the hot sections of an aero-engine turbine are subjected to high thermal and mechanical stresses in the aggressive atmosphere of high-velocity combustion gases. They have therefore to exhibit a high mechanical strength at elevated temperatures, but also a sufficient level of oxidation and corrosion resistance. Indeed, even if the use of a protective coating system on the component can be considered as a barrier against high
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Fig. 1. Effect of substitution solutes on strength of chromium (980–1320 °C) [1].
temperature oxidation and corrosion processes, the alloy itself has to exhibit a good environmental resistance to avoid catastrophic damage in case of erosion or complete removal of this coating. Pure chromium exhibits poor mechanical strength. Unalloyed chromium shows 0.2% compressive yield strength of about 200 MPa at room temperature and about 50 MPa at 1200 °C [9]. Various strengthening modes have therefore been previously investigated including solid solution strengthening, oxide dispersion strengthening (ODS), and precipitation strengthening by carbide, boride, and Laves phase particles. Solid solution strengthening of chromium by addition of refractory elements such as Ta, Nb, V, Re, W, Mo [1] or Ru [10] can lead to substantial increases of high-temperature strength (Figs. 1 and 2). However, an adverse effect of this solid solution strengthening effect is sometimes to decrease the low-temperature ductility. Some alloying elements such as Re, Ru, Fe and Co do not exhibit this drawback [1], but addition of Ru or Re increases the density that may be a serious handicap for aero-engine
applications [3]. Addition of oxide particles such as ThO2 [16,17], MgO [11,8] and Y2O3 [19,20] in chromium was shown to improve its room-temperature ductility [11], to retard subscale Cr2N formation [18] or to lower the DBTT [16]. Veigel [16] showed an oxide dispersion strengthening (ODS) effect at low temperature (150–200 °C), but this beneficial effect unfortunately vanishes at high temperature (1093 °C). Moreover, the techniques used to produce these ODS materials are not easily implemented (powder metallurgy þconsolidation or vapour-deposition þ consolidation). Efficient high-temperature strengthening of Cr was obtained by precipitation of carbides (Ti, Nb, Hf, Zr or Ta monocarbides) or/and borides (Zr, (Nb,Cr), (Ta,Cr) or Cr borides) but the DBTT is significantly increased as compared to pure chromium (between 177 and 371 °C) [21–23]. These alloys have however the advantage to be prepared by conventional methods (induction melting). Strengthening by Cr2X Laves phase particles (X ¼Ta, Hf, Nb, Zr or Ti) was tested in various systems [4–7,24,25]. These Cr–Cr2X alloys generally show some potential for high-temperature strength but have limited low-temperature fracture toughness. Moreover, the Cr2X Laves phase is considered to be less-oxidation resistant than the Cr solid solution matrix. Lastly, the Cr2X Laves phases are prone to nitriding to a much greater depth than the Cr matrix in Cr–Cr2Nb and Cr–Cr2Ta alloys [5]. As precipitation of the ordered γ′-Ni3Al type intermetallic compound is the most efficient way to strengthen the nickel-based superalloys, we have explored the possibility to use a similar strengthening mode in chromium-based alloys. Alloys of this study were therefore designed in order to induce precipitation of strengthening coherent or semi-coherent phase particles in a chromium-based solid solution, while promoting the formation of a protective external oxide layer. The precipitates must have a crystalline structure derived from the A2 structure of α-Cr and a lattice parameter close to that of Cr to ensure some coherency between the matrix and the precipitates such as to mimic the γ/γ′ nickel-based superalloys. The lattice parameter of chromium being 0.28839 nm [26], only compounds with a lattice parameter in the 0.284–0.293 nm range have been selected in order to keep the lattice mismatch within the range7 1.6% at room temperature. The lattice mismatch δ is defined as:
δ = 2(aCr –a precipitate )/(aCr + a precipitate )
Fig. 2. Compressive 0.2% yield strength vs. temperature for pure chromium, binary chromium alloys and TMS-75 alloy (a nickel-based single crystal superalloy) [9].
323
(1)
where aCr and aprecipitate are the lattice parameters of α-Cr and the precipitate phase, respectively. On the other hand, in order to get a good oxidation resistance at high temperature, one objective was to make the alloys prone to the formation of a protective external oxide scale. To ensure a sufficient protection at high temperature, this oxide scale must fulfil several criteria: (i) thermodynamic stability, (ii) slow growth kinetics (i.e. low diffusion coefficient of oxidizing species within the oxide layer), (iii) good adhesion to the substrate under thermal cyclic conditions, (iv) minimum thermal expansion mismatch with the substrate alloy, (v) limited reaction with the atmosphere, (vi) efficient barrier against nitrogen penetration. In practice, chromia (Cr2O3) and alumina (Al2O3) are among the most protective oxides likely to form on high temperature Fe-, Co- or Ni-based alloys. Of course, chromium-based alloys preferentially develop stable chromia external scales when exposed in air at high temperature, but the resulting protective effect is limited to temperatures below 900 °C. Indeed, at higher temperatures, Cr2O3 reacts with oxygen to form volatile oxide species (CrO3) [27]. This reaction is even more significant in presence of high oxygen pressure and high velocity gases. On the other side, alumina which grows preferentially on high-aluminium content Ni-based superalloys is
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stable at temperatures above 900 °C. Moreover, transport mechanisms are slower in alumina than in chromia which implies slower diffusion-controlled growth rate of the oxide layer and longer lifetimes. Finally, alumina is more thermodynamically stable than chromia [28]. Another drawback of chromia is the precipitation of nitrides beneath the oxide scale [5,7,29], indicating, as suggested by Klopp [1] and confirmed by Zheng and Young [30], that a chromia scale is permeable to nitrogen. On the contrary, as demonstrated by Han et al. [31] by studying the oxidation behaviour of Ni–Cr–Al alloys, alumina scales prevent nitrogen to reach the metal and ensure an effective protection. Considering both high temperature oxidation and mechanical resistance criteria therefore led to the choice of aluminium as a major alloying element to promote precipitation of intermetallic coherent strengthening particles and growth of a protective external alumina scale. It was thus chosen to develop chromiumbased alloys strengthened by precipitation of β′-Ni2AlTi phase resulting from additions of significant amounts of nickel, aluminium and titanium. The Ni2AlTi phase has a cubic ordered L21 structure (also known as Heusler phase) and its lattice parameter is equal to 0.5850 nm [32] leading to a lattice mismatch of 1.42% (using Eq. (1) with doubled α-Cr lattice parameter to take into account the 8 unit cells of the Heusler structure). Three quaternary alloys were thus defined (Table 1). Aluminium is in excess as compared with titanium in order to obtain alumina-forming alloys.
3. Experimental procedures Small ingots (30 to 50 cm3) of each alloy were prepared either by arc melting in water-cooled copper crucible or by cold crucible levitation induction melting of high-purity materials under argon atmosphere. Some alloys were also prepared by vacuum induction melting in a ceramic mould under argon atmosphere and die–cast into a steel mould to obtain rods with a diameter of 15 mm (30 cm3). Densities of samples were calculated by dividing their mass determined using high-precision weighing by their volume measured using helium pycnometry. All the ageing or pre-oxidation heat treatments were conducted in horizontal tube furnaces in argon atmosphere. The sample temperature was measured through a type S thermocouple placed above and close to it. Once the heat treatment duration reached, the sample is cooled down. Air cooling (AC-generally used) consists in quick withdrawal of the sample from the furnace. For salt water quenching, the sample is rapidly removed from the furnace and plunged into the aqueous media. For slow 5 K/min. cooling, the sample is kept in the furnace which cooling rate is managed through the temperature control unit. The alloy microstructures were examined using an Olympus PMG3 optical microscope (OM), a Zeiss DSM 962 scanning electron microscope (SEM) attached with an energy dispersive X-ray spectrometer (EDS) and a JEOL 200CX Transmission Electron Microscope (TEM) operating at 200 kV. The thin foils for TEM observation were prepared by mechanical polishing of 3 mm diameter discs, then by twin jet electro-polishing using a solution of Table 1 Nominal chemical compositions (at%) of the chromium-based alloys. Alloy
Cr
Ni
Al
Ti
RT7 RT6 RT4
80 70 60
10 15 20
6 9 12
4 6 8
35 vol% of 2-butylethanol, 59 vol% of methanol and 6 vol% of perchloric acid at 20 °C and 20 V (Struers Tenupol 3). Some thin foils were also prepared by ion beam polishing (Gatan 691 Precision Ion Polishing System) at 5 kV and 0.5 mA. X-Ray diffraction profiles were recorded using a Philips PW 1380 X-ray diffractometer with CuKα radiation selected thanks to a graphite monochromator. Cyclic oxidation tests on polished samples (Fig. 3a) were performed at 1100 °C in a tube furnace open to laboratory air. The samples were hanged up in the furnace and underwent one-hour thermal cycles (Fig. 3b). Every hour, the samples were automatically removed from the furnace and air fan cooled. Each sample was initially measured and weighted. It was periodically weighted in order to assess the rate of oxide scale growth via the mass variation during the cyclic oxidation test. Before cutting of cross-sections for microstructural analyses, electroless copper plating and nickel plating were carried out on oxidized samples to prevent damage of the oxide scales.
4. Experimental results 4.1. Density The densities of the experimental alloys were determined in order to make comparisons with typical blade nickel-based superalloys (Table 2). These chromium-based alloys offer a 11–13% mass reduction when compared with conventionally cast IN100 superalloy (one of the lightest blade superalloy) or a 23–25% mass reduction if compared with a third-generation single-crystal superalloy such as René N6. 4.2. Microstructure OM observations of the as-cast microstructures revealed dendrites with a light contrast surrounded by darker interdendritic phases (Fig. 4). Higher magnification observations (Fig. 5) allowed distinguishing interdendritic areas with a massive single phase, and other ones with an eutectic microstructure. The fraction of interdendritic phases increases as the Cr content decreases (Fig. 4). Three analysis methods were used in order to identify the different phases. EDS analyses of the dendrites indicate a very high content of Cr, while single phase interdendritic areas exhibit chemical compositions compatible with that of the β′-Ni2AlTi phase (Table 3). The X-ray diffraction diagram of the RT4 alloy is shown in Fig. 6. Typical diffraction peaks of the β′-Ni2AlTi compound confirm the presence of this phase. The measured lattice parameters are respectively 0.289 nm for α-Cr phase and 0.586 nm for β′-Ni2AlTi phase. The calculated lattice mismatch of about – 1.59% corresponds to the upper limit of the range defined in § 2. It's worth noting that the interdendritic β′-Ni2AlTi phase of RT4 alloy contains 7.6 at% Cr (Table 3). This result agrees well with those experimentally obtained by Doğan et al. in two Cr–Ni–Al–Ti alloys [33]. These authors confirmed this result by first-principles calculations which predict that Cr prefers substitution for the Ni site over Al or Ti sites in the Ni2AlTi phase. SEM observations of etched samples (Murakami reagent) revealed very fine particles (mean diameter of 50 nm) within the Crrich dendrites (Fig. 7). TEM observations performed on RT7 alloy thin foils confirmed these observations (Fig. 8a). On some bright field images, depending on the diffraction conditions, the presence of these particles was evidenced by the typical “coffee-grain” strain contrast of quasi-coherent particles (Fig. 8b) where two closely spaced parallel strain-lobes surround a line of no contrast [34]. Such lobes of contrast are characteristic of relatively large and misfitting precipitates [35]. Analysis of selected area electron
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Fig. 3. Cyclic oxidation: (a) sample geometry and (b) temperature variation during an oxidation cycle. Table 2 Densities of chromium-based alloys and typical blade nickel-based superalloys. Alloy
RT4
RT6
RT7
IN100
René N6
Density (g cm 3)
6.74
6.80
6.87
7.75
8.97
diffraction patterns (SADPs) obtained for a 〈011〉 zone axis, coupled with diffraction pattern simulation, confirmed the two-phase microstructure of the dendrites with a α-Cr matrix and quasi-coherent β′-Ni2AlTi precipitates (Fig. 9). Similar SADPs were observed by Doğan et al. in a Cr–5Ni–4.5Al–0.5Ti (at%) alloy after solutioning at 1400 °C and aging at 1000 °C with a subsequent furnace cool [36]. According to Doğan, the characteristic L21 diffraction spots originate from small amount of Ni2AlTi domains present in the coarse B2–NiAl precipitates identified in this alloy beside a finer population of NiAl nanoprecipitates. In order to evaluate the influence of the β′-Ni2AlTi precipitates on the mechanical strength of the alloys, some room temperature microhardness tests were carried out on RT4 alloy samples being applied various ageing heat treatments. Some resulting microstructures are illustrated by the SEM micrographs of Fig. 10. Ageing for one hour at 1000 °C and 1200 °C induces a relatively strong coarsening of the β′-Ni2AlTi precipitates as compared with ageing for one hour at 800 °C. The microhardness values are compared in Fig. 11. The precipitation strengthening is maximum after ageing at 800 °C, higher than for the as-cast alloy. Ageing at 1000 °C and 1200 °C decreases the alloy strength as compared with the as-cast microstructure. At 1200 °C, the cooling rate has little effect on the microhardness data.
4.3. Environmental resistance 4.3.1. Cyclic oxidation tests at 1100 °C Cyclic oxidation tests were performed on the three chromiumbased alloys at 1100 °C in air (300 one-hour cycles). Mass variations are compared in Fig. 12 to results obtained with the same experimental conditions on AM1 and MC-NG nickel-based superalloy samples. AM1 alloy [37] is a first-generation industrial single crystal blade superalloy (used in the Snecma M88 engine) while MC-NG is a fourth-generation single crystal blade superalloy under development [38]. The chromium-based alloys exhibit a continuous mass increase during the 150 first cycles and afterwards some of these alloys begin to degrade. A very weak mass gain is observed for MC-NG superalloy when AM1 superalloy experiences, after a slight mass gain, a continuous loss of mass. SEM observations performed on cross-sections of AM1 and MCNG samples showed a thin (o 5 mm) external oxide scale after 1000 oxidation cycles at 1100 °C. EDS analyses indicated that these oxide scales are made of Al2O3 (Fig. 13). The higher cyclic oxidation resistance of the MC-NG alloy is due to a denser, more flat and adherent layer of Al2O3 than for AM1 alloy. This difference could be attributed, at least partly, to the intentional additions of 0.1 wt% hafnium and 0.1 wt% silicon in MC-NG alloy [39]. Some TiN particles were observed beneath the surface in AM1 alloy because of discontinuities in the Al2O3 scale. For the chromium-based alloys, no continuous Al2O3 layer is observed on cross-sections (Fig. 14). Typically, the alloy area modified by the air exposure includes three distinct zones: – the external zone (zone A) is made of chromia containing fine alumina particles (30 to 50 mm thickness),
Fig. 4. As-cast microstructures – (a) RT4 alloy, (b) RT6 alloy and (c) RT7 alloy (OM images).
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Fig. 5. As-cast microstructures – (a) RT4 alloy and (b) RT6 alloy (OM images). Table 3 Results of the EDS analysis of the dendrite and interdendritic areas in RT4 alloy (at%). Phase
Cr
Ni
Al
Ti
Dendrite (α-Crþ β′-Ni2AlTi precipitates) Interdendritic massive phase (β′-Ni2AlTi)
89.3 7.6
4.1 45.8
4.6 28.1
2.0 18.5
– the internal zone (zone C) where chromia particles are observed in the chromium dendrites. The depth of the whole zone affected by the air exposure is about 100 mm after 300 cycles at 1100 °C. These observations show that formation of a chromia scale does not prevent further diffusion of oxygen and nitrogen deeply into the chromium-based alloys. It could be also pointed out that despite its high aluminium content (12 at%), RT4 alloy still developed a non-protective chromia scale. The mass variation Δm of oxidation samples can generally be described by a parabolic curve expressed by the following relation:
⎛ Δm ⎞2 ⎜ ⎟ = k t p ⎝ S ⎠
Fig. 6. X-ray diffraction diagram of as-cast RT4 alloy.
– the intermediate zone (zone B) also contains chromia particles and alumina particles (edgings in RT7 alloy) or chromia particles and TiN or AlN particles (in RT4 alloy). In RT7 alloy, TiN particles decorate some α/β′ boundaries,
(2)
where S is the initial sample surface, t is the time and kp is the parabolic rate constant for oxidation (g2 cm 4 s 1). When the Cr2O3 oxide scale is submitted to high temperatures (T o900 °C), formation of volatile oxide species (CrO3) has to be taken into account. Actually, the reaction Cr2O3 (solid) þ 3/2 O2 (gas) 2 2 CrO3 (solid) has to be considered during oxidation of chromium at high temperature and under high oxygen pressures. This reaction can result in significant scale thinning by evaporation [27]. To take into account this phenomenon, Stott [27] proposes the following model:
kp dy = − kv dt y
Fig. 7. β′-Ni2AlTi precipitates in the α-Cr phase: (a) RT4 alloy and (b) RT7 alloy (SEM images).
(3)
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Fig. 8. β′-Ni2AlTi precipitates in the α-Cr phase of RT7 alloy (bright field TEM images).
Fig. 9. (a) Actual and (b) simulated [101] zone axis diffraction patterns of the α-Cr matrix and β′-Ni2AlTi precipitates in the RT7 alloy.
where y is the oxide scale thickness (proportional to Δm/S), kp is the parabolic rate constant describing the diffusion process in the scale and kv is the constant describing the rate of volatilization. Thus, for the alloys developing a Cr2O3 oxide scale, the parabolic rate constant kp must be calculated using the modified relation:
⎛ Δm ⎞2 ⎜ + k vt ⎟ = k pt ⎝ S ⎠
(4)
At 1100 °C, the kv value is estimated to be 6 10 9 g cm 2 s 1[40]. The kp values calculated from the experimental results and using Eq. (4) are within the range 1.7–5.2 10 10 g2 cm 4 s 1. These values are well above the expected value for an aluminaformer alloy which should not exceed 5 10 12 g2 cm 4 s 1 at 1100 °C according to [41]. No relationship can be established between the cyclic oxidation kinetics and the aluminium concentration. 4.3.2. Solutions to enhance alumina formation 4.3.2.1. High aluminium content alloys. One solution envisaged to promote alumina scale formation on the chromium-base alloys was to increase their aluminium concentration. Two additional compositions were thus designed with higher aluminium contents. Nominal compositions of these alloys are presented in Table 4 together with their measured densities. RT14 alloy derives from RT4 alloy by increasing the Al/Ni and Al/Ti ratios in order to
have aluminium in excess and promote alumina formation during oxidation. RT15 alloy was designed to increase the amount of Ni2TiAl phase. The higher fraction of aluminium-rich phases was also expected to promote alumina formation. These alloys were prepared by vacuum induction melting in a ceramic mould under argon. Cyclic oxidation tests performed at 1100 °C did not evidence any significant differences concerning the oxidation behaviour of RT14 and RT15 alloys as compared with RT4 and RT7 (Fig. 15). SEM observations and EDS analyses of oxidized samples revealed that RT14 and RT15 alloys are still chromia-formers despite their higher aluminium contents. However, SEM observations showed a thin alumina scale above some large interdendritic areas (Fig. 16). Under this alumina scale, no oxidation or nitriding phenomenon is locally observed, that confirms the protective nature of this oxide. 4.3.2.2. Pre-oxidation at 1300 °C. One other solution tested to increase the high temperature oxidation resistance of chromiumbase alloys was to form an external alumina layer on these alloys before cyclic oxidation at 1100 °C. Indeed, a preliminary study showed that an isothermal exposure in air at 1380 °C led to the formation of such a protective alumina layer. Further work was therefore conducted on the quaternary RT4 alloy selected because of its high aluminium amount (12 at%) and prepared by arc melting. A pre-oxidation treatment was carried out at a lower temperature of 1300 °C during four hours to avoid incipient melting. Comparative cyclic oxidation tests at 1100 °C with 50 one-
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Fig. 10. Evolution of the Ni2AlTi precipitation in the α-Cr phase of RT4 alloy: (a) after 800 °C/1 h AC, (b) after 1000 °C/1 h AC and (c) after 1200 °C/1 h SWQ (AC: air cooling, CR: cooling rate, and SWQ: salt water quenching) (SEM images).
700 650
HV 1kg 20°C
600 550 500 450 400 350 300
as-cast
800°C/1h AC
800°C/24h AC
1000°C/1h AC
1000°C/24h AC
1200°C/1h CR 5K/min
1200°C/1h SWQ
Heat treatment Fig. 11. Room temperature microhardness of RT4 alloy for various ageing heat treatments (AC: air cooling, CR: cooling rate, and SWQ: salt water quenching).
Fig. 12. One-hour cyclic oxidation of chromium-based alloys and nickel-based superalloys at 1100 °C.
hour cycles were carried out on as-cast and pre-oxidized samples. At the end of the test, a smaller mass gain was measured for the pre-oxidized sample as compared with the as-cast one (Fig. 17). SEM observations and EDS analyses showed that a continuous layer of Cr2O3 was formed on both samples. However, the chromia scale thickness was about 27 mm on the as-cast sample compared to only 17 mm for the pre-oxidized sample. Furthermore, in the pre-oxidized sample, a thin discontinuous alumina layer formed under the chromia scale when these events were rare in the ascast sample. As for the high aluminium content alloys, a pre-oxidation at 1300 °C during four hours can therefore promote the formation of alumina but not as a continuous scale.
4.3.2.3. Surface aluminium enrichment. A third solution experimented to promote formation of a protective alumina scale during oxidation test was to preliminary enrich the sample surface with aluminium using a vapour-phase aluminizing process. Experiments were conducted on both quaternary RT4 and RT7 alloys. The vapour-phase aluminizing process was carried out during five hours at 1100 °C in hydrogen atmosphere. As an example, the micrography of Fig. 18 illustrates the case of the aluminized RT7 sample. An external Al-enriched area was observed on cross-section of the sample, appearing as a 20 mm thick layer darker than the bulk material in backscattered electron detection (BSE) mode. EDS analyses confirmed an aluminium enrichment in this zone where the aluminium concentration increased both in the Cr
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Fig. 13. Cross-section of cyclic oxidation sample of MC-NG nickel-based superalloy after 1000 one-hour cycles at 1100 °C (SEM image). Fig. 15. One-hour cyclic oxidation at 1100 °C of quaternary chromium-based alloys.
Fig. 16. Cross-section of cyclic oxidation sample of RT14 chromium-based alloy after 50 one-hour cycles at 1100 °C (SEM image).
Fig. 14. Cross-section of cyclic oxidation sample of RT7 chromium-based alloy after 300 one-hour cycles at 1100 °C (SEM image).
Table 4 Nominal chemical compositions (at%) and measured densities of the Al-enriched chromium-based alloys. Alloy
Cr
Ni
Al
Ti
Density (g cm 3)
RT14 RT15
60 50
20 25
15 15
5 10
6.71 6.64
dendrites and in the interdendritic phase as compared with the bulk sample (Fig. 18). Cyclic oxidation tests were performed at 1100 °C on RT4 samples as this alloy is richer in Al than RT7 that was supposed to be more favourable to formation of alumina. Samples were first aluminized during five hours at 1100 °C. After aluminization, one sample was subjected to a pre-oxidation treatment at 1300 °C during four hours. The specific weight changes after 50 one-hour cycles are compared in Fig. 17 with the results obtained on as-cast and pre-oxidized RT4 samples, without preliminary aluminization. The aluminizing treatment alone led to a lower mass gain compared to the as-cast material. However, the best result is obtained thanks to the combination of the aluminizing and the pre-oxidation treatments that leads to a dramatic decrease of the oxide
Fig. 17. One-hour cyclic oxidation at 1100 °C of RT4 alloy with various pretreatments.
growth rate, the mass gain being strongly reduced after 50 oxidation cycles. Optical micrographs of the cross-section of the aluminized sample after 50 oxidation cycles showed that the oxide scale formed is permeable to oxygen and nitrogen. Microstructural damages spread to a depth of about 100 mm (Fig. 19a). On the contrary, no such damages were observed on the cross-section of the aluminized then pre-oxidized sample (Fig. 19b). SEM observations and EDS analyses indicated that the oxide scale formed on the surface of this sample is made of an external oxide layer of TiO2 and of an internal oxide layer of Al2O3 (thicknesso10 mm) (Fig. 20). 4.4. High-temperature creep resistance High-temperature compression creep tests were carried out on
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Fig. 18. Cross-section and local EDS analyses of aluminized sample of RT7 chromium-based alloy (BSE-SEM image).
specimens of as-cast RT4 and RT7 alloys and, for the sake of comparison, on specimens of reference materials: IN100 conventionally cast (CC) nickel-based superalloy and CMSX-2 and MC2 〈001〉 single crystal (SC) nickel-based superalloys, all strengthened by L12 Ni3Al-based γ′ precipitates, and 〈001〉 single crystal specimens of a β-NiAl-based alloy strengthened by precipitation of β′-Ni2AlTi Heusler phase. The dimensions of the parallelepipedic specimens were 3 3 6 mm3. The creep tests were conducted at 850 and 1100 °C by incremental step-loading. Minimum creep rates were measured after each stress increment and stabilization of the creep rate and are compared in Figs. 21 and 22. For the chromium-based alloys, the creep tests were performed under dynamic argon to prevent excessive environmental damages. Some data of the literature are also reported for pure chromium [42], for a TiAl-based alloy (Ti–47Al–2Cr–2Nb, in at%) [43] and for a niobium silicide-based alloy (Nb–8Hf–25Ti–2Al– 2Cr–16Si, in at%) [44]. Creep tests are compressive ones except for pure chromium, TiAl-based alloy and MC2 superalloy. Atmosphere and temperature creep tests are specified for each material in Figs. 21 and 22. At 850 and 1100 °C, the minimum creep rate of the RT7 alloy is lower than that of RT4. At 850 °C, the creep resistances of the chromium-based alloys RT4 and RT7 are much higher than that of pure chromium at 816 °C (Fig. 21). CC IN100 and SC CMSX-2 superalloys are more creep resistant than RT4 and RT7 alloys, but the chromium-based alloys exhibit minimum creep rates comparable to those of TiAlbased or NiAl (β/β′) alloys. At 1100 °C, the minimum creep rate of RT7 alloy is about 100 times lower than the tensile creep rate of pure chromium [42] at 982 °C (Fig. 22). But comparison with the niobium silicide-based alloy and with the SC CMSX-2 superalloy crept in vacuum or tested in air shows that RT4 and RT7 alloys present lower creep
resistances at this high temperature.
5. Concluding remarks Introduction of significant amounts of Al, Ni and Ti in experimental Cr-based alloys has demonstrated the possibility to induce precipitation of the intermetallic β′-Ni2AlTi phase in a α-Cr matrix. As-cast alloys prepared by arc melting or vacuum induction melting exhibit a dendritic microstructure with α-Cr-based dendrites containing fine dispersions of β′-Ni2AlTi precipitates and interdendritic areas containing massive β′ phase or α/β′ eutectic phase. The fraction of interdendritic phases increases with the total amount of β′-forming elements, Al, Ni and Ti. The β′ precipitates are quasi-coherent with the α-Cr matrix with a lattice mismatch of – 1.59%. Room temperature micro-hardness tests show that ageing heat treatment can control the size of the β′ precipitates and influence the strength of the alloys. Compressive creep tests performed at 850 °C and 1100 °C evidence creep rates significantly lower for the quaternary alloys as compared with pure chromium. Even if we cannot totally exclude a solid solution strengthening effect of the Cr-based matrix, this dramatic effect is likely due to a β′ precipitate strengthening effect. It must be pointed out that the RT4 alloy is less creep-resistant than the RT7 alloy although its total content of Al, Ni and Ti alloying elements is higher. This can be explained by a fraction of interdendritic phases higher in RT4 than in RT7 and a fraction of α/β′ dendrites correlatively higher. Thus the total content of β′-former elements could still be reduced to minimize the amount of interdendritic phases while keeping a high strength α/β′ microstructure within the dendrites in order to enhance the creep resistance.
Fig. 19. Cross-section of cyclic oxidation samples of RT4 alloy after 50 one-hour cycles at 1100 °C: (a) aluminized sample and (b) aluminized and pre-oxidized 4 h at 1300 °C (OM images).
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Fig. 20. Cross-section of the aluminized and pre-oxidized sample of RT4 alloy after 50 one-hour cycles at 1100 °C (SEM images).
Fig. 21. Minimum creep rate at 850 °C of RT4 and RT7 alloys compared to hightemperature materials.
the formation of a continuous and dense Al2O3 surface layer which protects the alloy from further degradation by lower temperature oxidation or nitridation mechanisms. Cyclic oxidation tests performed at 1100 °C demonstrated the huge beneficial effect of this method on the oxidation behaviour of the quaternary Cr-based alloys. A dramatic reduction of the mass gain rate is obtained as compared with as-cast, pre-oxidized alloys without prior aluminization, or aluminized alloys without further pre-oxidation. These preliminary results on new Cr-based quaternary alloys strengthened by Ni2AlTi precipitation pave the way for further development and investigations aiming at evaluating the real potential of such alloys as structural materials for high temperature components in gas turbine engines. Optimization of the chemistry and of the microstructure of these alloys seems still possible while keeping in mind the necessity to protect them from environmental aggressions. Finally, the powder metallurgy route (pre-alloyed powders) could be considered for the production of these materials with enhanced microstructure. The expected reduction of the chemical heterogeneity and the resulting refined microstructure due to the rapid solidification step is assumed to be a way to reach improved mechanical and environmental resistance.
Acknowledgements The authors would like to thank S. Navéos for vapour-phase aluminizing processing. They also acknowledge T. Thiounn, V. Mercerolle and J. Gantier for their contribution to this study. Fig. 22. Minimum creep rate at 1100 °C of RT4 and RT7 alloys compared to hightemperature materials.
The high temperature oxidation and nitridation resistance of Cr-based alloys is a key issue as their natural Cr2O3 oxide scale is volatile above 900 °C. Despite the high Al contents of these alloys, oxidation at 1100 °C leads to the formation of a chromia scale which does not prevent further diffusion of oxygen and nitrogen into the chromium-based alloys. However, an efficient way to develop an Al2O3 protective oxide scale at high temperature on the quaternary alloys was therefore identified. An aluminization process followed by a pre-oxydation treatment at 1300 °C promotes
References [1] W.D. Klopp, in: C.T. Sims, W.C. Hagel (Eds.), The Superalloys, John Wiley & Sons Inc., New-York, USA, 1972, pp. 175–196. [2] W.D. Klopp, J. Less Common Met. 42 (1975) 261–278. [3] Y.F. Gu, H. Harada, Y. Ro, JOM 42 (2004) 28–33. [4] C.T. Liu, P.F. Tortorelli, J.A. Horton, C.A. Carmichael, Mater. Sci. Eng. A214 (1996) 23–32. [5] M.P. Brady, J.H. Zhu, C.T. Liu, P.F. Tortorelli, L.R. Walker, Intermetallics 8 (2000) 1111–1118. [6] R.H. Tien, J.H. Zhu, C.T. Liu, L.R. Walker, Intermetallics 13 (2005) 361–366. [7] M.P. Brady, C.T. Liu, J.H. Zhu, P.F. Tortorelli, L.R. Walker, Scr. Mater. 52 (2005) 815–819. [8] Y.F. Gu, Y. Ro, H. Harada, Metall. Mater. Trans. A 35A (2004) 3329–3331.
332
D. Locq et al. / Materials Science & Engineering A 647 (2015) 322–332
[9] Y. Ro, Y. Koizumi, S. Nakazawa, T. Kobayashi, E. Bannai, H. Harada, Scr. Mater. 46 (2002) 331–335. [10] Y.F. Gu, H. Harada, Y. Ro, T. Kobayashi, Metall. Mater. Trans. A 36A (2005) 577–582. [11] G. Kunschert, W. Glatz, M. Janousek, R. Zach, M.P. Brady, I.G. Wright, in: Proceedings of the 16th International Plansee Seminar 2005: Powder Metallurgical High Performance Materials, vol. 1, 2005, pp. 1047–1059. [12] D. Locq, P. Caron, C. Ramusat, R. Mévrel, Adv. Mater. Res. 278 (2011) 569–574. [13] R. Wadsack, R. Pippan, B. Schedler, J. Nucl. Mater. 307–311 (2002) 701–704. [14] U. Holzwarth, H. Stamm, J. Nucl. Mater. 300 (2002) 161–177. [15] L. Royer, X. Ledoux, S. Mathieu, P. Steinmetz, Oxid. Met. 74 (2010) 79–92. [16] N.D. Veigel, et al., Development of a Chromium-Thoria Alloy, NASA CR-72901, 1971. [17] B.A. Wilcox, N.D. Veigel, A.H. Clauer, Met. Trans. 3 (1972) 273–283. [18] M.P. Brady, P. Sachenko, Scr. Mater. 52 (2005) 809–814. [19] PLANSEE, Ducropur, Ducrolloy, Highly Corrosion Resistant Materials, Technical Note, 〈www.plansee.com/〉, 2005. [20] H.-P. Martinz, W. Köck, T. Sakaki, J. Phys. IV 3 (1993) 205–213. [21] A.M. Filippi, Met. Trans. 3 (1972) 1727–1733. [22] A.M. Filippi, J. Less Common Met. 30 (1973) 153–158. [23] A.M. Filippi, Met. Trans. 5 (1974) 1423–1427. [24] K.S. Kumar, D.B. Miracle, Intermetallics 2 (1994) 257–274. [25] T. Takasugi, K.S. Kumar, C.T. Liu, E.H. Lee, Mater. Sci. Eng. A260 (1999) 108–123. [26] JCPDS File Number 06-0694, JCPDS International Center for Diffraction Data. [27] F.H. Stott, Rep. Prog. Phys. 50 (1987) 861–913. [28] B. Gleeson, in: R.A. Cottis, et al., (Eds.), Shreir's Corrosion, vol. 1, Elsevier B.V., Amsterdam, The Netherlands, 2010, pp. 180–194. [29] K. Taneichi, T. Narushima, Y. Iguchi, C. Ouchi, Mater. Trans. 47 (10) (2006)
[30] [31] [32] [33] [34] [35]
[36] [37] [38] [39]
[40] [41] [42] [43] [44]
2540–2546. X.G. Zheng, D.J. Young, Mater. Sci. Forum 251–254 (1997) 567–574. S. Han, D.J. Young, Oxid. Met. 55 (3–4) (2001) 223–242. JCPDS File Number 19-0034, JCPDS International Center for Diffraction Data. Ö.N. Doğan, X. Song, S. Chen, M.C. Gao, Intermetallics 35 (2013) 33–40. A.W. Wilson, J.M. Howe, A. Garg, R.D. Noebe, Mater. Sci. Eng. A289 (2000) 162–171. J.W. Edington, Interpretation of transmission electron micrographs, 3rd Volume of Monographs in Practical Electron Microscopy in Materials Science, Philips Technical Library, Eindhoven, Netherlands, 1975. Ö.N. Doğan, X. Song, D. Palacio, M.C. Gao, J. Mater. Sci. 49 (2014) 805–810. J.H. Davidson, A. Fredholm, T. Khan, J.-M. Théret, French Patent No. 2 557 598, 1983. P. Caron, in: T.M. Pollock, et al., (Eds.), Superalloys 2000, TMS, Warrendale, USA, 2000, pp. 737–746. P. Caron, S. Navéos, T. Khan, in: D. Coutsouradis, et al., (Eds.), Materials for Advanced Power Engineering Part II, Kluwer Academic Publishers, The Netherlands, 1994, pp. 1185–1194. K.P. Lillerud, P. Kofstad, J. Electrochem. Soc. 127 (1980) 2397–2419. H. Hindam, D.P. Wittle, Oxid. Met. 18 (5–6) (1982) 245–283. J.R. Stephens, W.D. Klopp, High-temperature creep of polycrystalline chromium, Technical Memorandum NASA TM X-2499, 1972. M. Thomas, S. Naka, in: Y.-W. Kim, D.M. Dimiduk, M.H. Loretto (Eds.), Gamma Titanium Aluminides, TMS, Warrendale, USA, 1999, pp. 633–640. B.P. Bewlay, M.R. Jackson, in: A. Kelly, B. Zweben (Eds.), Comprehensive Composite Materials, vol. 3, Elsevier, Amsterdam, The Netherlands, 2000, pp. 579–615.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Quaternary chromium-based alloys strengthened by Heusler phase precipitation D. Locq n, P. Caron, C. Ramusat, R. Mévrel Onera – The French Aerospace Lab, F-92322 Châtillon, France
art ic l e i nf o Article history: Received 29 July 2015 Received in revised form 7 September 2015 Accepted 8 September 2015 Available online 11 September 2015 Keywords: Chromium alloy Heusler-phase precipitation Electron microscopy Microanalysis Oxidation Creep
a b s t r a c t Five quaternary chromium-based alloys have been designed in order to precipitate β′-Ni2AlTi Heusler phase strengthening particles while forming a protective alumina scale. These alloys were prepared by arc or induction melting as small polycrystalline ingots. β′-Ni2AlTi Heusler phase was observed as massive particles in the interdendritic areas of the cast alloys and as fine precipitates within the dendrites. The level of aluminium added in these alloys was however not enough to promote the formation of a continuous alumina protective scale before or during cyclic oxidation tests at 1100 °C. Surface aluminium enrichment of cyclic oxidation samples was carried out by vapour-phase aluminization to promote formation of a protective alumina scale at 1100 °C. That resulted in a dramatic improvement of the cyclic oxidation resistance of the Cr-based alloys. Compressive creep test performed at 850 °C and 1100 °C demonstrated the significant strengthening effect of the β′-Ni2AlTi Heusler phase precipitates when compared with pure chromium. & 2015 Elsevier B.V. All rights reserved.
1. Introduction Nickel-based superalloys are used for decades as materials for high temperature aero-engine components such as blades, vanes, disks and casings. Development of superalloys with improved performance is becoming more and more difficult and costly and a new challenge for aero-engine manufacturers is to try to replace some of them with less dense and/or more mechanical and environmental resistant materials. This pressing goal is motivated by economical and performance requirements. Chromium-based alloys have been considered as potential candidates for high-temperature applications since the late 1940's because of their high melting point, high elastic modulus, low density and of the good oxidation resistance of these chromiaforming materials [1–3]. During the 1950's and 1960's, extensive research projects were mainly conducted in Australia, in the United States of America and in the Soviet Union [1]. However, these works highlighted or confirmed that chromium-based alloys could present some drawbacks, such as a high ductile-brittle transition temperature (DBTT) and a high nitrogen solubility which can cause severe embrittlement after exposure at high temperature in air. Moreover, it has not been proved till today that very good high-temperature mechanical properties could be reached. Given n Correspondence to: Onera, DMSM, BP72, 29 avenue de la Division Leclerc, 92322 Châtillon Cedex, France. Fax: þ33 146734164. E-mail address: [email protected] (D. Locq).
http://dx.doi.org/10.1016/j.msea.2015.09.033 0921-5093/& 2015 Elsevier B.V. All rights reserved.
these difficulties and the concomitant, rapid and successful development of nickel-based superalloys, effort to develop high temperature chromium-based alloys ceased in the late 1970's. However, research programs on such alloys have recently resumed. The potential applications are for general high temperature structural materials [3–8], or more specifically gas turbines [9–12], for nuclear plants [13,14] as well as for glass tools [15]. The present study aims at designing new chromium-based alloys with improved high-temperature mechanical and environmental resistance in order to propose an alternative to nickelbased superalloys for high temperature component applications in gas turbine engines, such as turbine vanes and blades. The alloys investigated here are strengthened by precipitates of the intermetallic β′-Ni2AlTi compound and are designed for developing a protective alumina scale instead of the chromia scale which usually forms on chromium-based alloys but which is unstable in the high temperature regime (T4 900 °C).
2. Alloy design Components in the hot sections of an aero-engine turbine are subjected to high thermal and mechanical stresses in the aggressive atmosphere of high-velocity combustion gases. They have therefore to exhibit a high mechanical strength at elevated temperatures, but also a sufficient level of oxidation and corrosion resistance. Indeed, even if the use of a protective coating system on the component can be considered as a barrier against high
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Fig. 1. Effect of substitution solutes on strength of chromium (980–1320 °C) [1].
temperature oxidation and corrosion processes, the alloy itself has to exhibit a good environmental resistance to avoid catastrophic damage in case of erosion or complete removal of this coating. Pure chromium exhibits poor mechanical strength. Unalloyed chromium shows 0.2% compressive yield strength of about 200 MPa at room temperature and about 50 MPa at 1200 °C [9]. Various strengthening modes have therefore been previously investigated including solid solution strengthening, oxide dispersion strengthening (ODS), and precipitation strengthening by carbide, boride, and Laves phase particles. Solid solution strengthening of chromium by addition of refractory elements such as Ta, Nb, V, Re, W, Mo [1] or Ru [10] can lead to substantial increases of high-temperature strength (Figs. 1 and 2). However, an adverse effect of this solid solution strengthening effect is sometimes to decrease the low-temperature ductility. Some alloying elements such as Re, Ru, Fe and Co do not exhibit this drawback [1], but addition of Ru or Re increases the density that may be a serious handicap for aero-engine
applications [3]. Addition of oxide particles such as ThO2 [16,17], MgO [11,8] and Y2O3 [19,20] in chromium was shown to improve its room-temperature ductility [11], to retard subscale Cr2N formation [18] or to lower the DBTT [16]. Veigel [16] showed an oxide dispersion strengthening (ODS) effect at low temperature (150–200 °C), but this beneficial effect unfortunately vanishes at high temperature (1093 °C). Moreover, the techniques used to produce these ODS materials are not easily implemented (powder metallurgy þconsolidation or vapour-deposition þ consolidation). Efficient high-temperature strengthening of Cr was obtained by precipitation of carbides (Ti, Nb, Hf, Zr or Ta monocarbides) or/and borides (Zr, (Nb,Cr), (Ta,Cr) or Cr borides) but the DBTT is significantly increased as compared to pure chromium (between 177 and 371 °C) [21–23]. These alloys have however the advantage to be prepared by conventional methods (induction melting). Strengthening by Cr2X Laves phase particles (X ¼Ta, Hf, Nb, Zr or Ti) was tested in various systems [4–7,24,25]. These Cr–Cr2X alloys generally show some potential for high-temperature strength but have limited low-temperature fracture toughness. Moreover, the Cr2X Laves phase is considered to be less-oxidation resistant than the Cr solid solution matrix. Lastly, the Cr2X Laves phases are prone to nitriding to a much greater depth than the Cr matrix in Cr–Cr2Nb and Cr–Cr2Ta alloys [5]. As precipitation of the ordered γ′-Ni3Al type intermetallic compound is the most efficient way to strengthen the nickel-based superalloys, we have explored the possibility to use a similar strengthening mode in chromium-based alloys. Alloys of this study were therefore designed in order to induce precipitation of strengthening coherent or semi-coherent phase particles in a chromium-based solid solution, while promoting the formation of a protective external oxide layer. The precipitates must have a crystalline structure derived from the A2 structure of α-Cr and a lattice parameter close to that of Cr to ensure some coherency between the matrix and the precipitates such as to mimic the γ/γ′ nickel-based superalloys. The lattice parameter of chromium being 0.28839 nm [26], only compounds with a lattice parameter in the 0.284–0.293 nm range have been selected in order to keep the lattice mismatch within the range7 1.6% at room temperature. The lattice mismatch δ is defined as:
δ = 2(aCr –a precipitate )/(aCr + a precipitate )
Fig. 2. Compressive 0.2% yield strength vs. temperature for pure chromium, binary chromium alloys and TMS-75 alloy (a nickel-based single crystal superalloy) [9].
323
(1)
where aCr and aprecipitate are the lattice parameters of α-Cr and the precipitate phase, respectively. On the other hand, in order to get a good oxidation resistance at high temperature, one objective was to make the alloys prone to the formation of a protective external oxide scale. To ensure a sufficient protection at high temperature, this oxide scale must fulfil several criteria: (i) thermodynamic stability, (ii) slow growth kinetics (i.e. low diffusion coefficient of oxidizing species within the oxide layer), (iii) good adhesion to the substrate under thermal cyclic conditions, (iv) minimum thermal expansion mismatch with the substrate alloy, (v) limited reaction with the atmosphere, (vi) efficient barrier against nitrogen penetration. In practice, chromia (Cr2O3) and alumina (Al2O3) are among the most protective oxides likely to form on high temperature Fe-, Co- or Ni-based alloys. Of course, chromium-based alloys preferentially develop stable chromia external scales when exposed in air at high temperature, but the resulting protective effect is limited to temperatures below 900 °C. Indeed, at higher temperatures, Cr2O3 reacts with oxygen to form volatile oxide species (CrO3) [27]. This reaction is even more significant in presence of high oxygen pressure and high velocity gases. On the other side, alumina which grows preferentially on high-aluminium content Ni-based superalloys is
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stable at temperatures above 900 °C. Moreover, transport mechanisms are slower in alumina than in chromia which implies slower diffusion-controlled growth rate of the oxide layer and longer lifetimes. Finally, alumina is more thermodynamically stable than chromia [28]. Another drawback of chromia is the precipitation of nitrides beneath the oxide scale [5,7,29], indicating, as suggested by Klopp [1] and confirmed by Zheng and Young [30], that a chromia scale is permeable to nitrogen. On the contrary, as demonstrated by Han et al. [31] by studying the oxidation behaviour of Ni–Cr–Al alloys, alumina scales prevent nitrogen to reach the metal and ensure an effective protection. Considering both high temperature oxidation and mechanical resistance criteria therefore led to the choice of aluminium as a major alloying element to promote precipitation of intermetallic coherent strengthening particles and growth of a protective external alumina scale. It was thus chosen to develop chromiumbased alloys strengthened by precipitation of β′-Ni2AlTi phase resulting from additions of significant amounts of nickel, aluminium and titanium. The Ni2AlTi phase has a cubic ordered L21 structure (also known as Heusler phase) and its lattice parameter is equal to 0.5850 nm [32] leading to a lattice mismatch of 1.42% (using Eq. (1) with doubled α-Cr lattice parameter to take into account the 8 unit cells of the Heusler structure). Three quaternary alloys were thus defined (Table 1). Aluminium is in excess as compared with titanium in order to obtain alumina-forming alloys.
3. Experimental procedures Small ingots (30 to 50 cm3) of each alloy were prepared either by arc melting in water-cooled copper crucible or by cold crucible levitation induction melting of high-purity materials under argon atmosphere. Some alloys were also prepared by vacuum induction melting in a ceramic mould under argon atmosphere and die–cast into a steel mould to obtain rods with a diameter of 15 mm (30 cm3). Densities of samples were calculated by dividing their mass determined using high-precision weighing by their volume measured using helium pycnometry. All the ageing or pre-oxidation heat treatments were conducted in horizontal tube furnaces in argon atmosphere. The sample temperature was measured through a type S thermocouple placed above and close to it. Once the heat treatment duration reached, the sample is cooled down. Air cooling (AC-generally used) consists in quick withdrawal of the sample from the furnace. For salt water quenching, the sample is rapidly removed from the furnace and plunged into the aqueous media. For slow 5 K/min. cooling, the sample is kept in the furnace which cooling rate is managed through the temperature control unit. The alloy microstructures were examined using an Olympus PMG3 optical microscope (OM), a Zeiss DSM 962 scanning electron microscope (SEM) attached with an energy dispersive X-ray spectrometer (EDS) and a JEOL 200CX Transmission Electron Microscope (TEM) operating at 200 kV. The thin foils for TEM observation were prepared by mechanical polishing of 3 mm diameter discs, then by twin jet electro-polishing using a solution of Table 1 Nominal chemical compositions (at%) of the chromium-based alloys. Alloy
Cr
Ni
Al
Ti
RT7 RT6 RT4
80 70 60
10 15 20
6 9 12
4 6 8
35 vol% of 2-butylethanol, 59 vol% of methanol and 6 vol% of perchloric acid at 20 °C and 20 V (Struers Tenupol 3). Some thin foils were also prepared by ion beam polishing (Gatan 691 Precision Ion Polishing System) at 5 kV and 0.5 mA. X-Ray diffraction profiles were recorded using a Philips PW 1380 X-ray diffractometer with CuKα radiation selected thanks to a graphite monochromator. Cyclic oxidation tests on polished samples (Fig. 3a) were performed at 1100 °C in a tube furnace open to laboratory air. The samples were hanged up in the furnace and underwent one-hour thermal cycles (Fig. 3b). Every hour, the samples were automatically removed from the furnace and air fan cooled. Each sample was initially measured and weighted. It was periodically weighted in order to assess the rate of oxide scale growth via the mass variation during the cyclic oxidation test. Before cutting of cross-sections for microstructural analyses, electroless copper plating and nickel plating were carried out on oxidized samples to prevent damage of the oxide scales.
4. Experimental results 4.1. Density The densities of the experimental alloys were determined in order to make comparisons with typical blade nickel-based superalloys (Table 2). These chromium-based alloys offer a 11–13% mass reduction when compared with conventionally cast IN100 superalloy (one of the lightest blade superalloy) or a 23–25% mass reduction if compared with a third-generation single-crystal superalloy such as René N6. 4.2. Microstructure OM observations of the as-cast microstructures revealed dendrites with a light contrast surrounded by darker interdendritic phases (Fig. 4). Higher magnification observations (Fig. 5) allowed distinguishing interdendritic areas with a massive single phase, and other ones with an eutectic microstructure. The fraction of interdendritic phases increases as the Cr content decreases (Fig. 4). Three analysis methods were used in order to identify the different phases. EDS analyses of the dendrites indicate a very high content of Cr, while single phase interdendritic areas exhibit chemical compositions compatible with that of the β′-Ni2AlTi phase (Table 3). The X-ray diffraction diagram of the RT4 alloy is shown in Fig. 6. Typical diffraction peaks of the β′-Ni2AlTi compound confirm the presence of this phase. The measured lattice parameters are respectively 0.289 nm for α-Cr phase and 0.586 nm for β′-Ni2AlTi phase. The calculated lattice mismatch of about – 1.59% corresponds to the upper limit of the range defined in § 2. It's worth noting that the interdendritic β′-Ni2AlTi phase of RT4 alloy contains 7.6 at% Cr (Table 3). This result agrees well with those experimentally obtained by Doğan et al. in two Cr–Ni–Al–Ti alloys [33]. These authors confirmed this result by first-principles calculations which predict that Cr prefers substitution for the Ni site over Al or Ti sites in the Ni2AlTi phase. SEM observations of etched samples (Murakami reagent) revealed very fine particles (mean diameter of 50 nm) within the Crrich dendrites (Fig. 7). TEM observations performed on RT7 alloy thin foils confirmed these observations (Fig. 8a). On some bright field images, depending on the diffraction conditions, the presence of these particles was evidenced by the typical “coffee-grain” strain contrast of quasi-coherent particles (Fig. 8b) where two closely spaced parallel strain-lobes surround a line of no contrast [34]. Such lobes of contrast are characteristic of relatively large and misfitting precipitates [35]. Analysis of selected area electron
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Fig. 3. Cyclic oxidation: (a) sample geometry and (b) temperature variation during an oxidation cycle. Table 2 Densities of chromium-based alloys and typical blade nickel-based superalloys. Alloy
RT4
RT6
RT7
IN100
René N6
Density (g cm 3)
6.74
6.80
6.87
7.75
8.97
diffraction patterns (SADPs) obtained for a 〈011〉 zone axis, coupled with diffraction pattern simulation, confirmed the two-phase microstructure of the dendrites with a α-Cr matrix and quasi-coherent β′-Ni2AlTi precipitates (Fig. 9). Similar SADPs were observed by Doğan et al. in a Cr–5Ni–4.5Al–0.5Ti (at%) alloy after solutioning at 1400 °C and aging at 1000 °C with a subsequent furnace cool [36]. According to Doğan, the characteristic L21 diffraction spots originate from small amount of Ni2AlTi domains present in the coarse B2–NiAl precipitates identified in this alloy beside a finer population of NiAl nanoprecipitates. In order to evaluate the influence of the β′-Ni2AlTi precipitates on the mechanical strength of the alloys, some room temperature microhardness tests were carried out on RT4 alloy samples being applied various ageing heat treatments. Some resulting microstructures are illustrated by the SEM micrographs of Fig. 10. Ageing for one hour at 1000 °C and 1200 °C induces a relatively strong coarsening of the β′-Ni2AlTi precipitates as compared with ageing for one hour at 800 °C. The microhardness values are compared in Fig. 11. The precipitation strengthening is maximum after ageing at 800 °C, higher than for the as-cast alloy. Ageing at 1000 °C and 1200 °C decreases the alloy strength as compared with the as-cast microstructure. At 1200 °C, the cooling rate has little effect on the microhardness data.
4.3. Environmental resistance 4.3.1. Cyclic oxidation tests at 1100 °C Cyclic oxidation tests were performed on the three chromiumbased alloys at 1100 °C in air (300 one-hour cycles). Mass variations are compared in Fig. 12 to results obtained with the same experimental conditions on AM1 and MC-NG nickel-based superalloy samples. AM1 alloy [37] is a first-generation industrial single crystal blade superalloy (used in the Snecma M88 engine) while MC-NG is a fourth-generation single crystal blade superalloy under development [38]. The chromium-based alloys exhibit a continuous mass increase during the 150 first cycles and afterwards some of these alloys begin to degrade. A very weak mass gain is observed for MC-NG superalloy when AM1 superalloy experiences, after a slight mass gain, a continuous loss of mass. SEM observations performed on cross-sections of AM1 and MCNG samples showed a thin (o 5 mm) external oxide scale after 1000 oxidation cycles at 1100 °C. EDS analyses indicated that these oxide scales are made of Al2O3 (Fig. 13). The higher cyclic oxidation resistance of the MC-NG alloy is due to a denser, more flat and adherent layer of Al2O3 than for AM1 alloy. This difference could be attributed, at least partly, to the intentional additions of 0.1 wt% hafnium and 0.1 wt% silicon in MC-NG alloy [39]. Some TiN particles were observed beneath the surface in AM1 alloy because of discontinuities in the Al2O3 scale. For the chromium-based alloys, no continuous Al2O3 layer is observed on cross-sections (Fig. 14). Typically, the alloy area modified by the air exposure includes three distinct zones: – the external zone (zone A) is made of chromia containing fine alumina particles (30 to 50 mm thickness),
Fig. 4. As-cast microstructures – (a) RT4 alloy, (b) RT6 alloy and (c) RT7 alloy (OM images).
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Fig. 5. As-cast microstructures – (a) RT4 alloy and (b) RT6 alloy (OM images). Table 3 Results of the EDS analysis of the dendrite and interdendritic areas in RT4 alloy (at%). Phase
Cr
Ni
Al
Ti
Dendrite (α-Crþ β′-Ni2AlTi precipitates) Interdendritic massive phase (β′-Ni2AlTi)
89.3 7.6
4.1 45.8
4.6 28.1
2.0 18.5
– the internal zone (zone C) where chromia particles are observed in the chromium dendrites. The depth of the whole zone affected by the air exposure is about 100 mm after 300 cycles at 1100 °C. These observations show that formation of a chromia scale does not prevent further diffusion of oxygen and nitrogen deeply into the chromium-based alloys. It could be also pointed out that despite its high aluminium content (12 at%), RT4 alloy still developed a non-protective chromia scale. The mass variation Δm of oxidation samples can generally be described by a parabolic curve expressed by the following relation:
⎛ Δm ⎞2 ⎜ ⎟ = k t p ⎝ S ⎠
Fig. 6. X-ray diffraction diagram of as-cast RT4 alloy.
– the intermediate zone (zone B) also contains chromia particles and alumina particles (edgings in RT7 alloy) or chromia particles and TiN or AlN particles (in RT4 alloy). In RT7 alloy, TiN particles decorate some α/β′ boundaries,
(2)
where S is the initial sample surface, t is the time and kp is the parabolic rate constant for oxidation (g2 cm 4 s 1). When the Cr2O3 oxide scale is submitted to high temperatures (T o900 °C), formation of volatile oxide species (CrO3) has to be taken into account. Actually, the reaction Cr2O3 (solid) þ 3/2 O2 (gas) 2 2 CrO3 (solid) has to be considered during oxidation of chromium at high temperature and under high oxygen pressures. This reaction can result in significant scale thinning by evaporation [27]. To take into account this phenomenon, Stott [27] proposes the following model:
kp dy = − kv dt y
Fig. 7. β′-Ni2AlTi precipitates in the α-Cr phase: (a) RT4 alloy and (b) RT7 alloy (SEM images).
(3)
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Fig. 8. β′-Ni2AlTi precipitates in the α-Cr phase of RT7 alloy (bright field TEM images).
Fig. 9. (a) Actual and (b) simulated [101] zone axis diffraction patterns of the α-Cr matrix and β′-Ni2AlTi precipitates in the RT7 alloy.
where y is the oxide scale thickness (proportional to Δm/S), kp is the parabolic rate constant describing the diffusion process in the scale and kv is the constant describing the rate of volatilization. Thus, for the alloys developing a Cr2O3 oxide scale, the parabolic rate constant kp must be calculated using the modified relation:
⎛ Δm ⎞2 ⎜ + k vt ⎟ = k pt ⎝ S ⎠
(4)
At 1100 °C, the kv value is estimated to be 6 10 9 g cm 2 s 1[40]. The kp values calculated from the experimental results and using Eq. (4) are within the range 1.7–5.2 10 10 g2 cm 4 s 1. These values are well above the expected value for an aluminaformer alloy which should not exceed 5 10 12 g2 cm 4 s 1 at 1100 °C according to [41]. No relationship can be established between the cyclic oxidation kinetics and the aluminium concentration. 4.3.2. Solutions to enhance alumina formation 4.3.2.1. High aluminium content alloys. One solution envisaged to promote alumina scale formation on the chromium-base alloys was to increase their aluminium concentration. Two additional compositions were thus designed with higher aluminium contents. Nominal compositions of these alloys are presented in Table 4 together with their measured densities. RT14 alloy derives from RT4 alloy by increasing the Al/Ni and Al/Ti ratios in order to
have aluminium in excess and promote alumina formation during oxidation. RT15 alloy was designed to increase the amount of Ni2TiAl phase. The higher fraction of aluminium-rich phases was also expected to promote alumina formation. These alloys were prepared by vacuum induction melting in a ceramic mould under argon. Cyclic oxidation tests performed at 1100 °C did not evidence any significant differences concerning the oxidation behaviour of RT14 and RT15 alloys as compared with RT4 and RT7 (Fig. 15). SEM observations and EDS analyses of oxidized samples revealed that RT14 and RT15 alloys are still chromia-formers despite their higher aluminium contents. However, SEM observations showed a thin alumina scale above some large interdendritic areas (Fig. 16). Under this alumina scale, no oxidation or nitriding phenomenon is locally observed, that confirms the protective nature of this oxide. 4.3.2.2. Pre-oxidation at 1300 °C. One other solution tested to increase the high temperature oxidation resistance of chromiumbase alloys was to form an external alumina layer on these alloys before cyclic oxidation at 1100 °C. Indeed, a preliminary study showed that an isothermal exposure in air at 1380 °C led to the formation of such a protective alumina layer. Further work was therefore conducted on the quaternary RT4 alloy selected because of its high aluminium amount (12 at%) and prepared by arc melting. A pre-oxidation treatment was carried out at a lower temperature of 1300 °C during four hours to avoid incipient melting. Comparative cyclic oxidation tests at 1100 °C with 50 one-
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Fig. 10. Evolution of the Ni2AlTi precipitation in the α-Cr phase of RT4 alloy: (a) after 800 °C/1 h AC, (b) after 1000 °C/1 h AC and (c) after 1200 °C/1 h SWQ (AC: air cooling, CR: cooling rate, and SWQ: salt water quenching) (SEM images).
700 650
HV 1kg 20°C
600 550 500 450 400 350 300
as-cast
800°C/1h AC
800°C/24h AC
1000°C/1h AC
1000°C/24h AC
1200°C/1h CR 5K/min
1200°C/1h SWQ
Heat treatment Fig. 11. Room temperature microhardness of RT4 alloy for various ageing heat treatments (AC: air cooling, CR: cooling rate, and SWQ: salt water quenching).
Fig. 12. One-hour cyclic oxidation of chromium-based alloys and nickel-based superalloys at 1100 °C.
hour cycles were carried out on as-cast and pre-oxidized samples. At the end of the test, a smaller mass gain was measured for the pre-oxidized sample as compared with the as-cast one (Fig. 17). SEM observations and EDS analyses showed that a continuous layer of Cr2O3 was formed on both samples. However, the chromia scale thickness was about 27 mm on the as-cast sample compared to only 17 mm for the pre-oxidized sample. Furthermore, in the pre-oxidized sample, a thin discontinuous alumina layer formed under the chromia scale when these events were rare in the ascast sample. As for the high aluminium content alloys, a pre-oxidation at 1300 °C during four hours can therefore promote the formation of alumina but not as a continuous scale.
4.3.2.3. Surface aluminium enrichment. A third solution experimented to promote formation of a protective alumina scale during oxidation test was to preliminary enrich the sample surface with aluminium using a vapour-phase aluminizing process. Experiments were conducted on both quaternary RT4 and RT7 alloys. The vapour-phase aluminizing process was carried out during five hours at 1100 °C in hydrogen atmosphere. As an example, the micrography of Fig. 18 illustrates the case of the aluminized RT7 sample. An external Al-enriched area was observed on cross-section of the sample, appearing as a 20 mm thick layer darker than the bulk material in backscattered electron detection (BSE) mode. EDS analyses confirmed an aluminium enrichment in this zone where the aluminium concentration increased both in the Cr
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Fig. 13. Cross-section of cyclic oxidation sample of MC-NG nickel-based superalloy after 1000 one-hour cycles at 1100 °C (SEM image). Fig. 15. One-hour cyclic oxidation at 1100 °C of quaternary chromium-based alloys.
Fig. 16. Cross-section of cyclic oxidation sample of RT14 chromium-based alloy after 50 one-hour cycles at 1100 °C (SEM image).
Fig. 14. Cross-section of cyclic oxidation sample of RT7 chromium-based alloy after 300 one-hour cycles at 1100 °C (SEM image).
Table 4 Nominal chemical compositions (at%) and measured densities of the Al-enriched chromium-based alloys. Alloy
Cr
Ni
Al
Ti
Density (g cm 3)
RT14 RT15
60 50
20 25
15 15
5 10
6.71 6.64
dendrites and in the interdendritic phase as compared with the bulk sample (Fig. 18). Cyclic oxidation tests were performed at 1100 °C on RT4 samples as this alloy is richer in Al than RT7 that was supposed to be more favourable to formation of alumina. Samples were first aluminized during five hours at 1100 °C. After aluminization, one sample was subjected to a pre-oxidation treatment at 1300 °C during four hours. The specific weight changes after 50 one-hour cycles are compared in Fig. 17 with the results obtained on as-cast and pre-oxidized RT4 samples, without preliminary aluminization. The aluminizing treatment alone led to a lower mass gain compared to the as-cast material. However, the best result is obtained thanks to the combination of the aluminizing and the pre-oxidation treatments that leads to a dramatic decrease of the oxide
Fig. 17. One-hour cyclic oxidation at 1100 °C of RT4 alloy with various pretreatments.
growth rate, the mass gain being strongly reduced after 50 oxidation cycles. Optical micrographs of the cross-section of the aluminized sample after 50 oxidation cycles showed that the oxide scale formed is permeable to oxygen and nitrogen. Microstructural damages spread to a depth of about 100 mm (Fig. 19a). On the contrary, no such damages were observed on the cross-section of the aluminized then pre-oxidized sample (Fig. 19b). SEM observations and EDS analyses indicated that the oxide scale formed on the surface of this sample is made of an external oxide layer of TiO2 and of an internal oxide layer of Al2O3 (thicknesso10 mm) (Fig. 20). 4.4. High-temperature creep resistance High-temperature compression creep tests were carried out on
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Fig. 18. Cross-section and local EDS analyses of aluminized sample of RT7 chromium-based alloy (BSE-SEM image).
specimens of as-cast RT4 and RT7 alloys and, for the sake of comparison, on specimens of reference materials: IN100 conventionally cast (CC) nickel-based superalloy and CMSX-2 and MC2 〈001〉 single crystal (SC) nickel-based superalloys, all strengthened by L12 Ni3Al-based γ′ precipitates, and 〈001〉 single crystal specimens of a β-NiAl-based alloy strengthened by precipitation of β′-Ni2AlTi Heusler phase. The dimensions of the parallelepipedic specimens were 3 3 6 mm3. The creep tests were conducted at 850 and 1100 °C by incremental step-loading. Minimum creep rates were measured after each stress increment and stabilization of the creep rate and are compared in Figs. 21 and 22. For the chromium-based alloys, the creep tests were performed under dynamic argon to prevent excessive environmental damages. Some data of the literature are also reported for pure chromium [42], for a TiAl-based alloy (Ti–47Al–2Cr–2Nb, in at%) [43] and for a niobium silicide-based alloy (Nb–8Hf–25Ti–2Al– 2Cr–16Si, in at%) [44]. Creep tests are compressive ones except for pure chromium, TiAl-based alloy and MC2 superalloy. Atmosphere and temperature creep tests are specified for each material in Figs. 21 and 22. At 850 and 1100 °C, the minimum creep rate of the RT7 alloy is lower than that of RT4. At 850 °C, the creep resistances of the chromium-based alloys RT4 and RT7 are much higher than that of pure chromium at 816 °C (Fig. 21). CC IN100 and SC CMSX-2 superalloys are more creep resistant than RT4 and RT7 alloys, but the chromium-based alloys exhibit minimum creep rates comparable to those of TiAlbased or NiAl (β/β′) alloys. At 1100 °C, the minimum creep rate of RT7 alloy is about 100 times lower than the tensile creep rate of pure chromium [42] at 982 °C (Fig. 22). But comparison with the niobium silicide-based alloy and with the SC CMSX-2 superalloy crept in vacuum or tested in air shows that RT4 and RT7 alloys present lower creep
resistances at this high temperature.
5. Concluding remarks Introduction of significant amounts of Al, Ni and Ti in experimental Cr-based alloys has demonstrated the possibility to induce precipitation of the intermetallic β′-Ni2AlTi phase in a α-Cr matrix. As-cast alloys prepared by arc melting or vacuum induction melting exhibit a dendritic microstructure with α-Cr-based dendrites containing fine dispersions of β′-Ni2AlTi precipitates and interdendritic areas containing massive β′ phase or α/β′ eutectic phase. The fraction of interdendritic phases increases with the total amount of β′-forming elements, Al, Ni and Ti. The β′ precipitates are quasi-coherent with the α-Cr matrix with a lattice mismatch of – 1.59%. Room temperature micro-hardness tests show that ageing heat treatment can control the size of the β′ precipitates and influence the strength of the alloys. Compressive creep tests performed at 850 °C and 1100 °C evidence creep rates significantly lower for the quaternary alloys as compared with pure chromium. Even if we cannot totally exclude a solid solution strengthening effect of the Cr-based matrix, this dramatic effect is likely due to a β′ precipitate strengthening effect. It must be pointed out that the RT4 alloy is less creep-resistant than the RT7 alloy although its total content of Al, Ni and Ti alloying elements is higher. This can be explained by a fraction of interdendritic phases higher in RT4 than in RT7 and a fraction of α/β′ dendrites correlatively higher. Thus the total content of β′-former elements could still be reduced to minimize the amount of interdendritic phases while keeping a high strength α/β′ microstructure within the dendrites in order to enhance the creep resistance.
Fig. 19. Cross-section of cyclic oxidation samples of RT4 alloy after 50 one-hour cycles at 1100 °C: (a) aluminized sample and (b) aluminized and pre-oxidized 4 h at 1300 °C (OM images).
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Fig. 20. Cross-section of the aluminized and pre-oxidized sample of RT4 alloy after 50 one-hour cycles at 1100 °C (SEM images).
Fig. 21. Minimum creep rate at 850 °C of RT4 and RT7 alloys compared to hightemperature materials.
the formation of a continuous and dense Al2O3 surface layer which protects the alloy from further degradation by lower temperature oxidation or nitridation mechanisms. Cyclic oxidation tests performed at 1100 °C demonstrated the huge beneficial effect of this method on the oxidation behaviour of the quaternary Cr-based alloys. A dramatic reduction of the mass gain rate is obtained as compared with as-cast, pre-oxidized alloys without prior aluminization, or aluminized alloys without further pre-oxidation. These preliminary results on new Cr-based quaternary alloys strengthened by Ni2AlTi precipitation pave the way for further development and investigations aiming at evaluating the real potential of such alloys as structural materials for high temperature components in gas turbine engines. Optimization of the chemistry and of the microstructure of these alloys seems still possible while keeping in mind the necessity to protect them from environmental aggressions. Finally, the powder metallurgy route (pre-alloyed powders) could be considered for the production of these materials with enhanced microstructure. The expected reduction of the chemical heterogeneity and the resulting refined microstructure due to the rapid solidification step is assumed to be a way to reach improved mechanical and environmental resistance.
Acknowledgements The authors would like to thank S. Navéos for vapour-phase aluminizing processing. They also acknowledge T. Thiounn, V. Mercerolle and J. Gantier for their contribution to this study. Fig. 22. Minimum creep rate at 1100 °C of RT4 and RT7 alloys compared to hightemperature materials.
The high temperature oxidation and nitridation resistance of Cr-based alloys is a key issue as their natural Cr2O3 oxide scale is volatile above 900 °C. Despite the high Al contents of these alloys, oxidation at 1100 °C leads to the formation of a chromia scale which does not prevent further diffusion of oxygen and nitrogen into the chromium-based alloys. However, an efficient way to develop an Al2O3 protective oxide scale at high temperature on the quaternary alloys was therefore identified. An aluminization process followed by a pre-oxydation treatment at 1300 °C promotes
References [1] W.D. Klopp, in: C.T. Sims, W.C. Hagel (Eds.), The Superalloys, John Wiley & Sons Inc., New-York, USA, 1972, pp. 175–196. [2] W.D. Klopp, J. Less Common Met. 42 (1975) 261–278. [3] Y.F. Gu, H. Harada, Y. Ro, JOM 42 (2004) 28–33. [4] C.T. Liu, P.F. Tortorelli, J.A. Horton, C.A. Carmichael, Mater. Sci. Eng. A214 (1996) 23–32. [5] M.P. Brady, J.H. Zhu, C.T. Liu, P.F. Tortorelli, L.R. Walker, Intermetallics 8 (2000) 1111–1118. [6] R.H. Tien, J.H. Zhu, C.T. Liu, L.R. Walker, Intermetallics 13 (2005) 361–366. [7] M.P. Brady, C.T. Liu, J.H. Zhu, P.F. Tortorelli, L.R. Walker, Scr. Mater. 52 (2005) 815–819. [8] Y.F. Gu, Y. Ro, H. Harada, Metall. Mater. Trans. A 35A (2004) 3329–3331.
332
D. Locq et al. / Materials Science & Engineering A 647 (2015) 322–332
[9] Y. Ro, Y. Koizumi, S. Nakazawa, T. Kobayashi, E. Bannai, H. Harada, Scr. Mater. 46 (2002) 331–335. [10] Y.F. Gu, H. Harada, Y. Ro, T. Kobayashi, Metall. Mater. Trans. A 36A (2005) 577–582. [11] G. Kunschert, W. Glatz, M. Janousek, R. Zach, M.P. Brady, I.G. Wright, in: Proceedings of the 16th International Plansee Seminar 2005: Powder Metallurgical High Performance Materials, vol. 1, 2005, pp. 1047–1059. [12] D. Locq, P. Caron, C. Ramusat, R. Mévrel, Adv. Mater. Res. 278 (2011) 569–574. [13] R. Wadsack, R. Pippan, B. Schedler, J. Nucl. Mater. 307–311 (2002) 701–704. [14] U. Holzwarth, H. Stamm, J. Nucl. Mater. 300 (2002) 161–177. [15] L. Royer, X. Ledoux, S. Mathieu, P. Steinmetz, Oxid. Met. 74 (2010) 79–92. [16] N.D. Veigel, et al., Development of a Chromium-Thoria Alloy, NASA CR-72901, 1971. [17] B.A. Wilcox, N.D. Veigel, A.H. Clauer, Met. Trans. 3 (1972) 273–283. [18] M.P. Brady, P. Sachenko, Scr. Mater. 52 (2005) 809–814. [19] PLANSEE, Ducropur, Ducrolloy, Highly Corrosion Resistant Materials, Technical Note, 〈www.plansee.com/〉, 2005. [20] H.-P. Martinz, W. Köck, T. Sakaki, J. Phys. IV 3 (1993) 205–213. [21] A.M. Filippi, Met. Trans. 3 (1972) 1727–1733. [22] A.M. Filippi, J. Less Common Met. 30 (1973) 153–158. [23] A.M. Filippi, Met. Trans. 5 (1974) 1423–1427. [24] K.S. Kumar, D.B. Miracle, Intermetallics 2 (1994) 257–274. [25] T. Takasugi, K.S. Kumar, C.T. Liu, E.H. Lee, Mater. Sci. Eng. A260 (1999) 108–123. [26] JCPDS File Number 06-0694, JCPDS International Center for Diffraction Data. [27] F.H. Stott, Rep. Prog. Phys. 50 (1987) 861–913. [28] B. Gleeson, in: R.A. Cottis, et al., (Eds.), Shreir's Corrosion, vol. 1, Elsevier B.V., Amsterdam, The Netherlands, 2010, pp. 180–194. [29] K. Taneichi, T. Narushima, Y. Iguchi, C. Ouchi, Mater. Trans. 47 (10) (2006)
[30] [31] [32] [33] [34] [35]
[36] [37] [38] [39]
[40] [41] [42] [43] [44]
2540–2546. X.G. Zheng, D.J. Young, Mater. Sci. Forum 251–254 (1997) 567–574. S. Han, D.J. Young, Oxid. Met. 55 (3–4) (2001) 223–242. JCPDS File Number 19-0034, JCPDS International Center for Diffraction Data. Ö.N. Doğan, X. Song, S. Chen, M.C. Gao, Intermetallics 35 (2013) 33–40. A.W. Wilson, J.M. Howe, A. Garg, R.D. Noebe, Mater. Sci. Eng. A289 (2000) 162–171. J.W. Edington, Interpretation of transmission electron micrographs, 3rd Volume of Monographs in Practical Electron Microscopy in Materials Science, Philips Technical Library, Eindhoven, Netherlands, 1975. Ö.N. Doğan, X. Song, D. Palacio, M.C. Gao, J. Mater. Sci. 49 (2014) 805–810. J.H. Davidson, A. Fredholm, T. Khan, J.-M. Théret, French Patent No. 2 557 598, 1983. P. Caron, in: T.M. Pollock, et al., (Eds.), Superalloys 2000, TMS, Warrendale, USA, 2000, pp. 737–746. P. Caron, S. Navéos, T. Khan, in: D. Coutsouradis, et al., (Eds.), Materials for Advanced Power Engineering Part II, Kluwer Academic Publishers, The Netherlands, 1994, pp. 1185–1194. K.P. Lillerud, P. Kofstad, J. Electrochem. Soc. 127 (1980) 2397–2419. H. Hindam, D.P. Wittle, Oxid. Met. 18 (5–6) (1982) 245–283. J.R. Stephens, W.D. Klopp, High-temperature creep of polycrystalline chromium, Technical Memorandum NASA TM X-2499, 1972. M. Thomas, S. Naka, in: Y.-W. Kim, D.M. Dimiduk, M.H. Loretto (Eds.), Gamma Titanium Aluminides, TMS, Warrendale, USA, 1999, pp. 633–640. B.P. Bewlay, M.R. Jackson, in: A. Kelly, B. Zweben (Eds.), Comprehensive Composite Materials, vol. 3, Elsevier, Amsterdam, The Netherlands, 2000, pp. 579–615.