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Use of multi-step loading small punch test to investigate the ductile-to-brittle transition behaviour of a thermally sprayed CoNiCrAlY coating
Author’s Accepted Manuscript Use of multi-step loading small punch test to investigate the ductile-to-brittle transition behaviour of a thermally sprayed CoNiCrAlY coating H. Chen, T.H. Hyde www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(16)31320-X http://dx.doi.org/10.1016/j.msea.2016.10.097 MSA34299
To appear in: Materials Science & Engineering A Received date: 5 September 2016 Revised date: 25 October 2016 Accepted date: 26 October 2016 Cite this article as: H. Chen and T.H. Hyde, Use of multi-step loading small punch test to investigate the ductile-to-brittle transition behaviour of a thermally sprayed CoNiCrAlY coating, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.10.097 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Use of multi-step loading small punch test to investigate the ductile-to-brittle transition behaviour of a thermally sprayed CoNiCrAlY coating H. Chena*, T. H. Hydeb a
Department of Mechanical, Materials and Manufacturing Engineering, Faculty of Science
and Engineering, University of Nottingham Ningbo China, Ningbo 315100, China b
Department of Mechanical, Materials and Manufacturing Engineering, Faculty of
Engineering, University of Nottingham, Nottingham NG7 2RD, UK *
Corresponding Author. Tel.: +86-574-88180946; Fax: +86-574-88187462. E-mail address:
[email protected]
Abstract Small punch tests under multi-step loading conditions were performed to characterise the ductile-to-brittle transition behaviour of a high velocity oxy-fuel (HVOF) thermally sprayed CoNiCrAlY (Co-31.7%Ni-20.8%Cr-8.1%Al-0.5%Y (wt%)) coating. Small punch specimens, 8 mm in diameter and ~ 0.4 mm in thickness, were tested between 21 ºC (RT) and 600 ºC. A 100 N load was applied in increments every 30 mins as a step to investigate the coating deformation at different temperatures. The displacement and strain obtained from multi-step loading SPTs at each load increment were small below 500 °C but a significant increase was noted at 600 °C. The strain rate behaviour was more distinct at 600 °C and large plastic deformation were shown, which is likely due to the ductile-to-brittle transition occurred between 500 – 600 °C. Fractographic investigation revealed that the fracture surfaces at low temperatures exhibited flat features with isolated β-phase particles, indicating the inter-splat shearing and brittle failure, whereas the main fracture mode was dominated by extensive ductile tearing at 600 °C.
Keywords: Small punch test; MCrAlY coating; HVOF spraying; DBTT; Strain 1
1. Introduction During service, gas turbine blades are exposed to high thermal stresses, oxidation and hot corrosion. Thermal barrier coating (TBC) systems are widely used to protect superalloy substrates from these harsh operating environments [1-4]. A typical TBC system that deposited onto the superalloy substrates consists of a thermal insulating ceramic top coat, an aluminium-rich bond coat, which provides good adhesion and oxidation resistance, and a thermally grown oxide (TGO), predominantly alumina, forms at the interface between the bond coat and the top coat during service at elevated temperatures [5-9]. It is widely recognised that the slowly thickened TGO during service causes a progressive build-up of stress in the system, leading to spallation of the ceramic top coat [10-17]. Maintaining the mechanical integrity of the multi-layered TBC system is one of the major challenges in the coating lifetime [18]. Since the imposed stress and strain resulted from the TGO growth can be accommodated by the creep, yield and ductility characteristics of the bond coat, the mechanical behaviour of the bond coat are crucial to the durability of the overall system [1921]. The ability of the bond coat to accommodate the imposed stress and strain largely depends on the coating microstructure, in-service temperature and coating thickness [22]. Bond coats are typically either of the diffusion aluminide type or the overlay MCrAlY type (M = Co and/or Ni). Overlay MCrAlY coatings have become more widely used because of advantages such as lower cost, better control of composition and the possibility to employ complex alloys with tailored microstructures. MCrAlY alloys are typically complex multi-phase materials and can comprise, for example, fcc -Ni and bcc B2 β-NiAl phases [23]. The bcc β-NiAl phase, commonly present in MCrAlY coatings, is brittle at low temperatures and shows increasing ductility above a critical temperature termed the ductile-to-brittle transition temperature (DBTT) [24]. Hence, overlay MCrAlY coatings comprising both β-NiAl phase and the ductile
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fcc -Ni phase also exhibit a transition in strain to cracking across the DBTT. Cracking may occur in the coating due to the low temperature brittleness of the β-NiAl phase during turbine start-up or shut-down operations. Furthermore, the strain tolerance of the coating is affected by the coating thickness since thin MCrAlY coating behaves differently from a bulk alloy of the same chemistry at elevated temperatures [25]. It is therefore important to determine the mechanical characteristics and the DBTT of MCrAlY coatings with thicknesses as close as possible to those used in industrial TBC systems (typically ~200 µm thick in TBCs). To accurately determine the above coating properties, one method of testing which appears to have the potential to provide such measurements is the small punch test (SPT) [26]. The procedure of testing, geometrical dimensions of the set-up and required test specimen dimensions are specified in a European Code of Practice [27]. The SPT technique has been used for decades now and has proven to be a reliable tool for mechanical characterisation of materials when analysing small regions of structural components, such as heat affected zones of welded joints and thin coatings. The test under bi-axial stress state is carried out by applying a hemispherical punch to one surface of the specimen which is firmly clamped between two circular dies. Depending on the test configurations, it has been used for creep analysis [28-33] and characterisation of the mechanical properties, especially, the DBTT of various materials [34-37]. Hitherto, the SPT has been validated as an effective method for assessing the mechanical properties and DBTT of bulk steels used in power plants [38-40], irradiated steels [41, 42] and, more recently, coatings [43-46]. Kameda et al. have studied the local embrittlement effects and mechanical degradations of MCrAlY coatings [47, 48]. Eskner et al. have used this technique to evaluate the DBTT of a nickel aluminide coating to be around 760 °C [44]. Soltysiak et al. reported the SPT work on mechanical characterisation of a thermally sprayed nickel-based superalloy (IN625) [45]. However, using SPT to investigate the ductile-to-brittle transition behaviour of MCrAlY bond coats does not appear to have been previously reported. Therefore, in this paper, the results of
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small punch tests under multi-step loading conditions for a HVOF thermally sprayed freestanding CoNiCrAlY coating between 21 °C ~ 600 °C are presented. Stress and strain behaviour is examined for the purpose of qualitatively analysing the ductile-to-brittle transition behaviour of this coating. Microstructural failure investigations are conducted at various temperatures to identify the dominant failure mode.
2. Experimental 2.1. Materials The coatings used in the SPT experiments were prepared by HVOF thermal spraying using powder with the following nominal composition Co-31.7%Ni-20.8%Cr-8.1%Al-0.5%Y (wt%). The powder was obtained from Praxair (CO-210-24) and had a size range of 45+20 m with a chemically analysed oxygen content of 0.037 wt% O. The coatings were deposited onto mild steel substrates with dimensions 60 × 25 × 1.8 mm3 using a Met Jet III liquid fuel HVOF gun. The spraying procedure is detailed elsewhere [49]. The mild steel substrates were cleaned with ethanol and ground with 800 grade SiC paper in order to aid coating detachment after spraying. Coatings were sprayed to a thickness of approximately 0.5 mm and were then debonded from the mild steel substrates by bending around a mandrel to produce free-standing coatings. Detached coating samples were vacuum heat treated at 1100 °C for 2 h to reduce the porosity within the coating and to allow the secondary phase to precipitate. Specimens for the SPT in the form of 8 mm diameter discs, cut from the heat treated coatings by electro-discharge machining, were ground down from the as-deposited thickness and polished to a final thickness of approximately 430 m using 1 m diamond paste. The final thickness was controlled to within ± 8 m as measured by a digital micrometer.
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2.2. Multi-step loading SPT The SPTs were carried out in accordance with CEN Workshop Agreement, CWA 15627:2006 E [27], using the equipment shown schematically in Fig. 1. It comprises a steel frame, A, a sample holder, B, and a clamping head, C. A hemispherical punch, D, is set co-axially with the sample. The load is applied on the pan, E, and is transmitted to the sample via the vertical rod, F. Tests are conducted at elevated temperatures controlled by the furnace, G. The displacement at the centre of the disc, , was measured through the punch movements as a function of time and recorded by a LVDT transducer, H. Disc samples were set up at 21 ºC, 400 ºC, 450 ºC, 500 ºC and 600 ºC in air. At each temperature, a 100 N load was applied in increments every 30 mins as a step to investigate the increase in the displacement, strain and strain rate at each load increment. 450 ºC was previously reported as a temperature that stress relaxation started to occur for a similar MCrAlY alloy [50], therefore it is of interest to investigate the coating behaviour across this temperature.
Fig. 1. An image of the small punch test rig, the schematic diagram of the rig set-up and an enlarged view of the sample holder.
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To qualitatively characterise the equivalent stress and strain in SPT under multi-step loading conditions, the constant load (P) after each load increment can be converted to equivalent stress (σeq) using Eq. (1) [27]. Rs is the radius of the punch, which is 1 mm and Ks is a material dependent factor, taken to be 1 in this study. The equivalent strain (εeq) developed during tests can be related to the displacement (δ) of the punch using Eq. (2) [51]. The equivalent strain rate ( eq ) can be obtained by converting the displacement rate ( ̇ ) using Eq. (3), obtained by differentiating Eq. (2). It has to be emphasised that due to the nature of biaxial loading in SPT, the stress and strain characteristics are very complicated. Thus the values calculated from the above equations may not be the exact stress and strain experienced by the specimen, but only as the first approximation.
eq
P 3.33 K s a p0.2 Rs1.2 t
eq 0.15( )1.5 t
eq 0.225( ) ( ) t
1.5
(1) (2) (3)
2.3. Material characterisation Cross sections of the heat treated coating were mounted in a conductive resin, ground and diamond polished to a 1 m finish. Failed samples following SPT at different temperatures were mounted on carbon conductive tabs and the fracture surfaces were investigated at low and high magnifications to study the fracture behaviour. Prepared samples were examined in a FEI XL 30 scanning electron microscope (SEM) operated at 20 kV. Both backscattered electrons (BSE) and secondary electrons (SE) were used to form images and semi-quantitative energy dispersive X-ray analysis (EDX) was utilised to aid phase identification through chemical analysis. Electron backscattered diffraction (EBSD) was used to investigate the grain morphology of coating cross-sections using a Zeiss 1530 VP field emission gun
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scanning electron microscope (Carl Zeiss, Inc., Maple Grove, MN) with an EDAX Pegasus combined EBSD system (EDAX, Mahwah, NJ, USA). The EBSD patterns were recorded at a specimen tilt angle of 70 with an accelerating voltage of 20 kV and a beam current of 26 nA. EBSD maps were collected over an area of 50 m 50 m at a step size of 0.1 m. The EBSD data acquisition and processing were performed using the EDAX OIM5 software suite.
3. Results 3.1. Microstructure of the heat treated coatings The microstructures of the vacuum heat treated coatings are shown in Fig. 2. As previously reported, this CoNiCrAlY coating comprises two phases, the light contrast fcc γ-Ni phase and the dark contrast bcc β-(Ni,Co)Al phase. The volume fraction of the uniformly distributed β phase was found to be around 33% by image analysis. The thin, dark elongated features shown in Fig. 2(b) are oxide stringers that formed by oxidation of powders in-flight during the HVOF spraying process. EDX analysis showed them to be rich in oxygen and aluminium and they are presumably aluminium oxides with the volume fraction less than 1%. Fig. 2(c) shows a region of the coating imaged by EBSD and presented as an inversed pole figure map showing the grain morphology and orientation. Fig. 2(d) is the corresponding phase distribution map of the same region as Fig. 2(c) where red corresponds to β-phase precipitates and green represents the γ-phase matrix. Generally, the β-phase grains are monocrystalline and the γ-phase grains are mostly twinned. The random assortment of colours indicates that there is no preferred grain orientation for both phases. Finer grains were found around the particle boundaries, indicating areas of the coating which underwent melting and recrystallization during thermal spraying. The areas of larger grains highlight a powder particle which retained, in part, the original powder microstructure during HVOF thermal spraying. Both phases comprise grains which range in size from approximately 2 m to 0.5 m or even less. 7
Fig. 2. (a) Low magnification and (b) high magnification BSE micrographs showing the microstructures of the vacuum heat treated CoNiCrAlY coating prior to testing. (c) Inverse pole figure map and (d) phase distribution map derived from EBSD data showing the grain morphology and orientation of the coating. The β-phase is coloured red and the γ-phase is coloured green.
3.2. Results of multi-step loading SPT The multi-step loading curves in the form of displacement (δ) versus time for load increments of 100 N are shown in Fig. 3 at RT, 400 ºC, 450 ºC, 500 ºC and 600 ºC using 100 N as the load increment for each step. Discrete steps are seen at RT in Fig. 3(a), where the displacement increases by a small value at each load increment, denoted by the numbers 1, 2, 3, 4 in the plot, keeps constant before the next increment and eventually the coating failed after the final load increment at 400 N (step 4). The step loading curve at 400 ºC shown in Fig. 3(b) is similar to that at RT except that the specimen failed at the fifth load increment. A small displacement rate leading to fracture is found after the fourth load increment at 450 ºC in Fig. 3(c) and the sample failed when the load was increased to 500 N. Fig. 3(d) presents the 8
step loading curve at 500 ºC, small displacement rates are found to exist at the second and third load increment but become largely noticeable at the fourth load increment of 400 N. The displacement rates can be treated as the evidence of creep. This sample was left for creep to fracture at the fourth load increment when an obvious increase in the displacement with time was seen. At 600 ºC, Fig. 3(e), the sample has three load increments and the displacements at each load increment are two or three times larger than the ones at lower temperatures. It can also be observed that a positive displacement rate occurred after the first load increment of 100 N. A large displacement can be seen at the last load increment, which indicates extensive plastic deformation of the material at 600 ºC.
Fig. 3. Small punch multi-step loading curves from RT (21 ºC) to 600 ºC, (a) – (e), at the temperatures indicated in the plots. The numbers in the plots refer to the load step in increments of 100 N.
The instant displacements at each load increment for different temperatures are summarised in Fig. 4. By converting load to stress, displacement to strain and displacement rate to strain rate using Eq. (1) - (3), the corresponding comparison of strain and strain rate in log scale versus the imposed stresses can be obtained as shown in Fig. 5(a) and (b) respectively. It can be seen 9
from Fig. 5(a) specimens exhibit very small and similar strain values in the first three load increments between RT and 500 ºC. But the strain at the second and third load increment at 600 ºC are about one order of magnitude larger than those obtained at lower temperatures, which indicates that the deformation of the alloy increases dramatically. The test at 600 ºC finished after three steps whereas tests at other temperatures ended up with four or five steps. It is found that the strain obtained at 500 ºC and below lie in a good trend with stress increments in Fig. 5(a). The strain rate to stress behaviour below 500 ºC shown in Fig. 5(b) is very close to each other but a separation starts to occur for the test results at 500 ºC and more obvious at 600 ºC. The distinct behaviour in the strain rate at 600 ºC indicates the coating has undergone large plastic deformation.
Fig. 4. Plot of displacement versus load increments of 100 N from RT (21 ºC) to 600 ºC.
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Fig. 5. Plots of equivalent strain versus equivalent stress (a) and minimum strain rate versus equivalent stress (b).
3.3. Fracture surface investigation Fig. 6 shows the macroscopic fracture surfaces after multi-step loading tests at 21 ºC (a), 400 ºC (b), 500 ºC (c) and 600 ºC (d). The fracture surfaces shown at 21 ºC and 400 ºC did not exhibit macroscopic plastic deformation. Fast fracture was found when the final load increment was applied to the two specimens. The fracture surfaces at 500 ºC and 600 ºC show the evidence of cracking which initiated at the centre of the sample and propagated radially outwards to the clamping boundary. Evidence of macroscopic plastic deformation can be seen in Fig. 6(c) and (d).
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Fig. 6. Macroscopic fracture surfaces after SPTs at 21 C (a), 400 C (b) 500 C (c) and 600 C (d).
Fig. 7 depicts the microscopic characteristics of fracture surfaces. Areas of tearing are found from the fracture surface at 21 ºC in Fig. 7(a) due to the inherently ductile nature of γ matrix. Isolated angular β particles, which were proved to be Al rich by chemical analysis, are present on the fracture surface. Flat regions are visible in Fig. 7(a) and these distinctive features are probably the splat boundaries that formed due to the inter-splat shearing during testing. The fracture surface at 400 ºC still exhibits γ matrix “tearing” feature with discrete β particles in Fig. 7(b). Similarly at 500 ºC, γ matrix tearing remains the main deformation mechanism and the evidence of β particles on the fracture surface can also be found. When the temperature increased to 600 ºC, isolated β particles can hardly be observed and extensive tearing has become the dominant failure mode in Fig. 7(d).
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Fig. 7. Microscopic fracture surface characteristics of the CoNiCrAlY coating after multi-step loading SPT at 21 ºC (a), 400 ºC (b), 500 ºC (c) and 600 ºC (d), showing areas of tearing, isolated β particles and the transition of fracture behaviour.
4. Discussion 4.1. Data interpretation For the multi-step loading SPTs shown in Fig. 3, the displacement can be very small when the applied load, i.e. 100 N, is insufficient to produce enough deformation at low temperatures. Thus the percent error can be quite large when calculating the equivalent strain. Fig. 5(a) emphases this and is possibly the reason for the discrepancies of the two low strain points at 60 MPa. Apart from these two points, a general trend can be found, where strain is not very sensitive to temperature but varies with stress according to Eq. (4) when temperature is below 500 °C,
(4)
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It is generally found that the yield strength of MCrAlY coatings is about 500 MPa between 25 – 500 C and decreased to 200 MPa or less above 500 C [46]. Since the stresses applied in SPTs are much smaller than the yield strength at temperatures below 500 °C, the continuous plastic deformation did not occur after each load increment. The instant displacements of the last load increment are similar at 450 C and 500 C, implying that the ductility at both temperatures is close to each other. But it can be seen an increase in the strain rates at 500 °C in Fig. 5(b), which separates from the results at lower temperatures. This separation in the strain rate indicates the ductility of the coating starts to increase and the deformation has become thermally activated at 500 °C. Subsequently the step loading results at 600 C show a much larger increase in displacements and strain, implying that the stress-strain behaviour has become very sensitive to temperature. The strain rate behaviour becomes more distinct at 600 °C compared to that at lower temperatures in Fig. 5(b), exhibiting the further evidence for the ductile-to-brittle transition. It is indicated that the deformation has become plastic flow dominant and undergoes continuous plastic deformation at 600 °C from Fig. 5. According to the work reported by Noebe et al. [24], the yield strength of NiAl is about 250 MPa below 500 C and exhibits a sharp drop between 500 – 600 C. Thus, the stress applied at 600 C in multi-step loading SPTs in Fig. 5 can be very close or even higher to the yield strength. This could produce the extensive increase in displacement and strain, as seen in Fig. 5. Therefore, it can be concluded that the ductile-to-brittle transition of this CoNiCrAlY is likely to occur between 500-600 C, where the coating exhibits the brittle characteristics in deformation below 500 C but becomes prominently plastic deformable at 600 C.
4.2. Fracture mechanisms Since the microstructure of the coating shown in Fig. 2 consists of the γ-phase matrix and monocrystalline β-phase precipitates, it is possible that cracking occurs in the brittle β-phase grains at low temperatures and propagates in the ductile γ-phase matrix. The EBSD phase 14
map in Fig. 2(c) and (d) shows that regions of fine-grained β-phases around the particle boundary which can act as paths for crack initiation and are the reasons for the large flat features observed in Fig. 7(a). One of these paths can be referred to the periphery of the powder particle which has, in part, retained the original powder microstructure. Crack propagation along this path would leave a concave feature in the fracture surface that matches the shape of the powder particle due to the inter-splat shearing, such as the one shown in Fig. 7(a). At low temperatures, the energy required for crack propagation in the β-phase is less than the energy required for plastic flow. Since the volume fraction of the γ phase is about 70%, thus a mixture of γ-matrix tearing and isolated β particles are present, showing the β phases are still brittle. This is the reason why both brittle and ductile behaviour can be observed between 400-500 C in Fig. 7(b) and (c). Since the deformation of the β phase has been thermally activated at 600 C, thus less energy is required for the plastic flow than brittle cleavage in the β-phase. The gradual decrease in the threshold energy for plastic flow allows the β-phase to deform with the γ matrix, as shown in Fig. 7(d), causing the brittle to ductile transition to occur. The fracture surfaces show the tearing is the main failure mode above this DBTT, with a great similarity to the fracture surfaces reported at 750 C [31].
4.3. Comparison with previous studies The CoNiCrAlY coating alloy employed in this study consists of γ matrix phase and β-NiAl secondary phase. The γ matrix phase is ductile at all temperatures and only the β-NiAl phase exhibits the ductile to brittle transition behaviour. The temperature dependent ductility behaviour of β-NiAl reviewed by Noebe et al. [24] shows that the DBTT is between 500 – 700 C. It should be noted that the DBTT of MCrAlY depends on the amount of β-NiAl phase since only β-NiAl exhibits the ductile to brittle transition behaviour [18]. Meetham et al. [3] reported that diffusional aluminides which are predominantly brittle β phase exhibit higher DBTT than those overlay coatings with brittle β and ductile γ phase. Eskner et al. [44] found 15
that the DBTT of a nickel aluminide coating was about 760 C. Since the volume fraction of β-NiAl phase in the CoNiCrAlY coating used in this study is around 30%, the DBTT is expected to be less than 760 C. With regard to the DBTT of MCrAlY alloys, the available data are limited but values in the range 400-900 C are suggested compared to the present range 500-600 C, which shows good agreement among various MCrAlY alloys reviewed by Bose [18]. However, previous studies involve different MCrAlY compositions and phases, thus direct comparison with these studies are not applicable. The loading curves at 400 C and 450 C still exhibits brittle characteristics though the fracture surface indicates some matrix tearing exists due to the presence of ductile matrix. From the evolution of the fracture surfaces in Fig. 7, a clear difference between the fracture behaviour across the test temperature range can be seen, in which γ matrix tearing and isolated β particles were widely observed below 500 C but extensive tearing becomes dominant at 600 °C. Therefore, by combining the test results and fracture characteristics obtained from multi-step loading SPTs, it is believed that the brittle to ductile transition of this CoNiCrAlY coating occurs between 500 ~ 600 °C, where at 500 °C the transition is about to occur and 600 °C and above is clearly ductile tearing dominant.
5. Conclusions
In this work it is shown that the small punch multi-step loading test is a suitable method for qualitatively evaluating the ductile-to-brittle transition behaviour of the CoNiCrAlY bond coat materials.
The displacement and strain obtained at each load increment are small below 500 C, but the resultant strain at 600 C is about one order of magnitude larger. Distinct stress-strain behaviour is observed at 600 C, indicating that the coating has undergone extensive plastic deformation with an obvious increase in ductility.
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Evidence of inter-splat shearing, matrix tearing and isolated β particles were found at the fracture surfaces between RT and 500 C due to the presence of ductile γ-Ni matrix phase and brittle β-NiAl phase. Extensive ductile tearing becomes the dominant failure mode at 600 C due to the brittle to ductile transition occurs between 500 – 600 C.
Acknowledgements The authors would like to thank Mr. Shane Maskill for his skilled assistance in carrying out the small punch testing experiments at the University of Nottingham. The authors also thank Prof. Wei Sun and Prof. Graham McCartney for help discussions. The financial support from Provincial Applied Research Programme for Commonweal Technology funded by Science Technology Department of Zhejiang Province (No. 2016C31023) and Ningbo Natural Science Foundation Programme (No. 2016A610114) are acknowledged.
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[15] E.A.G. Shillington, D.R. Clarke, Acta Mater., 47 (1999) 1297-1305. [16] P.K. Wright, A.G. Evans, Curr. Opin. Solid State Mater. Sci., 4 (1999) 255-265. [17] D. Naumenko, V. Shemet, L. Singheiser, W. Quadakkers, J. Mater. Sci., 44 (2009) 16871703. [18] S. Bose, High Temperature Coatings, Elsevier Butterworth-Heinemann, Oxford, 2007. [19] K.J. Hemker, B.G. Mendis, C. Eberl, Mater. Sci. Eng. A, 483–484 (2008) 727-730. [20] U. Hermosilla, M.S.A. Karunaratne, I.A. Jones, T.H. Hyde, R.C. Thomson, Mater. Sci. Eng. A, 513–514 (2009) 302-310. [21] X. Gong, Y. Yang, Y. Ma, H. Peng, H. Guo, Mater. Sci. Eng. A, 673 (2016) 39-46. [22] A. Strang, E. Lang, Effect of Coatings on the Mechanical Properties of Superalloys, in: R. Brunetaud, D. Coutsouradis, T.B. Gibbons, Y. Lindblom, D.B. Meadowcroft, R. Stickler (Eds.) High Temperature Alloys for Gas Turbines, Springer, Netherlands, 1982, pp. 469-506. [23] J.R. Nicholls, JOM, 52 (2000) 28-35. [24] R.D. Noebe, R.R. Bowman, M.V. Nathal, Int. Mater. Rev., 38 (1993) 193-232. [25] M.P. Taylor, H.E. Evans, E.P. Busso, Z.Q. Qian, Acta Mater., 54 (2006) 3241-3252. [26] G.E. Lucas, MTA, 21 (1990) 1105-1119. [27] CEN Workshop Agreement. CWA 15627:2006 E, "Small punch test method for metallic materials," CEN, European Committee for Standardization, Brussels Belgium, 2006. [28] B.J. Kim, Y.B. Sim, J.H. Lee, M.K. Kim, B.S. Lim, Procedia Engineering, 10 (2011) 2579-2584. [29] S. Komazaki, T. Kato, Y. Kohno, H. Tanigawa, Mater. Sci. Eng. A, 510–511 (2009) 229233. [30] D.T. Blagoeva, R.C. Hurst, Mater. Sci. Eng. A, 510-511 (2009) 219-223. [31] H. Chen, T.H. Hyde, K.T. Voisey, D.G. McCartney, Mater. Sci. Eng. A, 585 (2013) 205213. [32] T.H. Hyde, M. Stoyanov, W. Sun, C.J. Hyde, J. Strain Anal. Eng. Des., 45 (2010) 141164. [33] H. Chen, T.H. Hyde, K.T. Voisey, D.G. McCartney, Proc IMechE Part L: J Materials: Design and Application, DOI 10.1177/1464420715622495 (2016). [34] S. Rasche, M. Kuna, Int. J. Pressure Vessels Piping, 125 (2015) 23-34. [35] T. Linse, M. Kuna, H.W. Viehrig, Mater. Sci. Eng. A, 614 (2014) 136-147. [36] K. Turba, R. Hurst, P. Hähner, Int. J. Pressure Vessels Piping, 111–112 (2013) 155-161. [37] S. Dryepondt, B.A. Pint, Surf. Coat. Technol., 205 (2010) 1195-1199. [38] F. Dobeš, K. Milička, Intermetallics, 18 (2010) 1357-1359. [39] J.H. Bulloch, Engineering Failure Analysis, 9 (2002) 511-534. [40] T.E. García, C. Rodríguez, F.J. Belzunce, C. Suárez, J. Alloys Compd., 582 (2014) 708717. [41] M.C. Kim, Y.J. Oh, B.S. Lee, Nucl. Eng. Des., 235 (2005) 1799-1805. [42] Z.X. Wang, H.J. Shi, J. Lu, P. Shi, X.F. Ma, Nucl. Eng. Des., 238 (2008) 3186-3193. [43] A.K. Ray, N. Roy, A. Kar, A.K. Ray, S.C. Bose, G. Das, J.K. Sahu, D.K. Das, B. Venkataraman, S.V. Joshi, Mater. Sci. Eng. A, 505 (2009) 96-104. [44] M. Eskner, R. Sandström, Surf. Coat. Technol., 165 (2003) 71-80. [45] S. Soltysiak, M. Selent, S. Roth, M. Abendroth, M. Hoffmann, H. Biermann, M. Kuna, Mater. Sci. Eng. A, 613 (2014) 259-263. [46] M. Eskner, R. Sandström, Surf. Coat. Technol., 200 (2006) 2695-2703. [47] J. Kameda, T.E. Bloomer, Y. Sugita, A. Ito, S. Sakurai, Mater. Sci. Eng. A, 229 (1997) 42-54. [48] J. Kameda, T.E. Bloomer, Y. Sugita, A. Ito, S. Sakurai, Mater. Sci. Eng. A, 234-236 (1997) 489-492. [49] S. Saeidi, K.T. Voisey, D.G. McCartney, J. Therm. Spray Technol., 20 (2011) 1-13.
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[50] A.A. Wereszczak, J.G. Hemrick, T.P. Kirkland, J.A. Haynes, T.J. Fitzgerald, J.E. Junkin, Proceedings of the International Gas Turbine and Aeroengine Congress and Exhibition, Stockholm, Sweden, (June, 1998) 98-GT-403. [51] X. Mao, M. Saito, H. Takahashi, Scr Metall Mater, 25 (1991) 2481-2485.
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PII: DOI: Reference:
S0921-5093(16)31320-X http://dx.doi.org/10.1016/j.msea.2016.10.097 MSA34299
To appear in: Materials Science & Engineering A Received date: 5 September 2016 Revised date: 25 October 2016 Accepted date: 26 October 2016 Cite this article as: H. Chen and T.H. Hyde, Use of multi-step loading small punch test to investigate the ductile-to-brittle transition behaviour of a thermally sprayed CoNiCrAlY coating, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.10.097 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Use of multi-step loading small punch test to investigate the ductile-to-brittle transition behaviour of a thermally sprayed CoNiCrAlY coating H. Chena*, T. H. Hydeb a
Department of Mechanical, Materials and Manufacturing Engineering, Faculty of Science
and Engineering, University of Nottingham Ningbo China, Ningbo 315100, China b
Department of Mechanical, Materials and Manufacturing Engineering, Faculty of
Engineering, University of Nottingham, Nottingham NG7 2RD, UK *
Corresponding Author. Tel.: +86-574-88180946; Fax: +86-574-88187462. E-mail address:
[email protected]
Abstract Small punch tests under multi-step loading conditions were performed to characterise the ductile-to-brittle transition behaviour of a high velocity oxy-fuel (HVOF) thermally sprayed CoNiCrAlY (Co-31.7%Ni-20.8%Cr-8.1%Al-0.5%Y (wt%)) coating. Small punch specimens, 8 mm in diameter and ~ 0.4 mm in thickness, were tested between 21 ºC (RT) and 600 ºC. A 100 N load was applied in increments every 30 mins as a step to investigate the coating deformation at different temperatures. The displacement and strain obtained from multi-step loading SPTs at each load increment were small below 500 °C but a significant increase was noted at 600 °C. The strain rate behaviour was more distinct at 600 °C and large plastic deformation were shown, which is likely due to the ductile-to-brittle transition occurred between 500 – 600 °C. Fractographic investigation revealed that the fracture surfaces at low temperatures exhibited flat features with isolated β-phase particles, indicating the inter-splat shearing and brittle failure, whereas the main fracture mode was dominated by extensive ductile tearing at 600 °C.
Keywords: Small punch test; MCrAlY coating; HVOF spraying; DBTT; Strain 1
1. Introduction During service, gas turbine blades are exposed to high thermal stresses, oxidation and hot corrosion. Thermal barrier coating (TBC) systems are widely used to protect superalloy substrates from these harsh operating environments [1-4]. A typical TBC system that deposited onto the superalloy substrates consists of a thermal insulating ceramic top coat, an aluminium-rich bond coat, which provides good adhesion and oxidation resistance, and a thermally grown oxide (TGO), predominantly alumina, forms at the interface between the bond coat and the top coat during service at elevated temperatures [5-9]. It is widely recognised that the slowly thickened TGO during service causes a progressive build-up of stress in the system, leading to spallation of the ceramic top coat [10-17]. Maintaining the mechanical integrity of the multi-layered TBC system is one of the major challenges in the coating lifetime [18]. Since the imposed stress and strain resulted from the TGO growth can be accommodated by the creep, yield and ductility characteristics of the bond coat, the mechanical behaviour of the bond coat are crucial to the durability of the overall system [1921]. The ability of the bond coat to accommodate the imposed stress and strain largely depends on the coating microstructure, in-service temperature and coating thickness [22]. Bond coats are typically either of the diffusion aluminide type or the overlay MCrAlY type (M = Co and/or Ni). Overlay MCrAlY coatings have become more widely used because of advantages such as lower cost, better control of composition and the possibility to employ complex alloys with tailored microstructures. MCrAlY alloys are typically complex multi-phase materials and can comprise, for example, fcc -Ni and bcc B2 β-NiAl phases [23]. The bcc β-NiAl phase, commonly present in MCrAlY coatings, is brittle at low temperatures and shows increasing ductility above a critical temperature termed the ductile-to-brittle transition temperature (DBTT) [24]. Hence, overlay MCrAlY coatings comprising both β-NiAl phase and the ductile
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fcc -Ni phase also exhibit a transition in strain to cracking across the DBTT. Cracking may occur in the coating due to the low temperature brittleness of the β-NiAl phase during turbine start-up or shut-down operations. Furthermore, the strain tolerance of the coating is affected by the coating thickness since thin MCrAlY coating behaves differently from a bulk alloy of the same chemistry at elevated temperatures [25]. It is therefore important to determine the mechanical characteristics and the DBTT of MCrAlY coatings with thicknesses as close as possible to those used in industrial TBC systems (typically ~200 µm thick in TBCs). To accurately determine the above coating properties, one method of testing which appears to have the potential to provide such measurements is the small punch test (SPT) [26]. The procedure of testing, geometrical dimensions of the set-up and required test specimen dimensions are specified in a European Code of Practice [27]. The SPT technique has been used for decades now and has proven to be a reliable tool for mechanical characterisation of materials when analysing small regions of structural components, such as heat affected zones of welded joints and thin coatings. The test under bi-axial stress state is carried out by applying a hemispherical punch to one surface of the specimen which is firmly clamped between two circular dies. Depending on the test configurations, it has been used for creep analysis [28-33] and characterisation of the mechanical properties, especially, the DBTT of various materials [34-37]. Hitherto, the SPT has been validated as an effective method for assessing the mechanical properties and DBTT of bulk steels used in power plants [38-40], irradiated steels [41, 42] and, more recently, coatings [43-46]. Kameda et al. have studied the local embrittlement effects and mechanical degradations of MCrAlY coatings [47, 48]. Eskner et al. have used this technique to evaluate the DBTT of a nickel aluminide coating to be around 760 °C [44]. Soltysiak et al. reported the SPT work on mechanical characterisation of a thermally sprayed nickel-based superalloy (IN625) [45]. However, using SPT to investigate the ductile-to-brittle transition behaviour of MCrAlY bond coats does not appear to have been previously reported. Therefore, in this paper, the results of
3
small punch tests under multi-step loading conditions for a HVOF thermally sprayed freestanding CoNiCrAlY coating between 21 °C ~ 600 °C are presented. Stress and strain behaviour is examined for the purpose of qualitatively analysing the ductile-to-brittle transition behaviour of this coating. Microstructural failure investigations are conducted at various temperatures to identify the dominant failure mode.
2. Experimental 2.1. Materials The coatings used in the SPT experiments were prepared by HVOF thermal spraying using powder with the following nominal composition Co-31.7%Ni-20.8%Cr-8.1%Al-0.5%Y (wt%). The powder was obtained from Praxair (CO-210-24) and had a size range of 45+20 m with a chemically analysed oxygen content of 0.037 wt% O. The coatings were deposited onto mild steel substrates with dimensions 60 × 25 × 1.8 mm3 using a Met Jet III liquid fuel HVOF gun. The spraying procedure is detailed elsewhere [49]. The mild steel substrates were cleaned with ethanol and ground with 800 grade SiC paper in order to aid coating detachment after spraying. Coatings were sprayed to a thickness of approximately 0.5 mm and were then debonded from the mild steel substrates by bending around a mandrel to produce free-standing coatings. Detached coating samples were vacuum heat treated at 1100 °C for 2 h to reduce the porosity within the coating and to allow the secondary phase to precipitate. Specimens for the SPT in the form of 8 mm diameter discs, cut from the heat treated coatings by electro-discharge machining, were ground down from the as-deposited thickness and polished to a final thickness of approximately 430 m using 1 m diamond paste. The final thickness was controlled to within ± 8 m as measured by a digital micrometer.
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2.2. Multi-step loading SPT The SPTs were carried out in accordance with CEN Workshop Agreement, CWA 15627:2006 E [27], using the equipment shown schematically in Fig. 1. It comprises a steel frame, A, a sample holder, B, and a clamping head, C. A hemispherical punch, D, is set co-axially with the sample. The load is applied on the pan, E, and is transmitted to the sample via the vertical rod, F. Tests are conducted at elevated temperatures controlled by the furnace, G. The displacement at the centre of the disc, , was measured through the punch movements as a function of time and recorded by a LVDT transducer, H. Disc samples were set up at 21 ºC, 400 ºC, 450 ºC, 500 ºC and 600 ºC in air. At each temperature, a 100 N load was applied in increments every 30 mins as a step to investigate the increase in the displacement, strain and strain rate at each load increment. 450 ºC was previously reported as a temperature that stress relaxation started to occur for a similar MCrAlY alloy [50], therefore it is of interest to investigate the coating behaviour across this temperature.
Fig. 1. An image of the small punch test rig, the schematic diagram of the rig set-up and an enlarged view of the sample holder.
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To qualitatively characterise the equivalent stress and strain in SPT under multi-step loading conditions, the constant load (P) after each load increment can be converted to equivalent stress (σeq) using Eq. (1) [27]. Rs is the radius of the punch, which is 1 mm and Ks is a material dependent factor, taken to be 1 in this study. The equivalent strain (εeq) developed during tests can be related to the displacement (δ) of the punch using Eq. (2) [51]. The equivalent strain rate ( eq ) can be obtained by converting the displacement rate ( ̇ ) using Eq. (3), obtained by differentiating Eq. (2). It has to be emphasised that due to the nature of biaxial loading in SPT, the stress and strain characteristics are very complicated. Thus the values calculated from the above equations may not be the exact stress and strain experienced by the specimen, but only as the first approximation.
eq
P 3.33 K s a p0.2 Rs1.2 t
eq 0.15( )1.5 t
eq 0.225( ) ( ) t
1.5
(1) (2) (3)
2.3. Material characterisation Cross sections of the heat treated coating were mounted in a conductive resin, ground and diamond polished to a 1 m finish. Failed samples following SPT at different temperatures were mounted on carbon conductive tabs and the fracture surfaces were investigated at low and high magnifications to study the fracture behaviour. Prepared samples were examined in a FEI XL 30 scanning electron microscope (SEM) operated at 20 kV. Both backscattered electrons (BSE) and secondary electrons (SE) were used to form images and semi-quantitative energy dispersive X-ray analysis (EDX) was utilised to aid phase identification through chemical analysis. Electron backscattered diffraction (EBSD) was used to investigate the grain morphology of coating cross-sections using a Zeiss 1530 VP field emission gun
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scanning electron microscope (Carl Zeiss, Inc., Maple Grove, MN) with an EDAX Pegasus combined EBSD system (EDAX, Mahwah, NJ, USA). The EBSD patterns were recorded at a specimen tilt angle of 70 with an accelerating voltage of 20 kV and a beam current of 26 nA. EBSD maps were collected over an area of 50 m 50 m at a step size of 0.1 m. The EBSD data acquisition and processing were performed using the EDAX OIM5 software suite.
3. Results 3.1. Microstructure of the heat treated coatings The microstructures of the vacuum heat treated coatings are shown in Fig. 2. As previously reported, this CoNiCrAlY coating comprises two phases, the light contrast fcc γ-Ni phase and the dark contrast bcc β-(Ni,Co)Al phase. The volume fraction of the uniformly distributed β phase was found to be around 33% by image analysis. The thin, dark elongated features shown in Fig. 2(b) are oxide stringers that formed by oxidation of powders in-flight during the HVOF spraying process. EDX analysis showed them to be rich in oxygen and aluminium and they are presumably aluminium oxides with the volume fraction less than 1%. Fig. 2(c) shows a region of the coating imaged by EBSD and presented as an inversed pole figure map showing the grain morphology and orientation. Fig. 2(d) is the corresponding phase distribution map of the same region as Fig. 2(c) where red corresponds to β-phase precipitates and green represents the γ-phase matrix. Generally, the β-phase grains are monocrystalline and the γ-phase grains are mostly twinned. The random assortment of colours indicates that there is no preferred grain orientation for both phases. Finer grains were found around the particle boundaries, indicating areas of the coating which underwent melting and recrystallization during thermal spraying. The areas of larger grains highlight a powder particle which retained, in part, the original powder microstructure during HVOF thermal spraying. Both phases comprise grains which range in size from approximately 2 m to 0.5 m or even less. 7
Fig. 2. (a) Low magnification and (b) high magnification BSE micrographs showing the microstructures of the vacuum heat treated CoNiCrAlY coating prior to testing. (c) Inverse pole figure map and (d) phase distribution map derived from EBSD data showing the grain morphology and orientation of the coating. The β-phase is coloured red and the γ-phase is coloured green.
3.2. Results of multi-step loading SPT The multi-step loading curves in the form of displacement (δ) versus time for load increments of 100 N are shown in Fig. 3 at RT, 400 ºC, 450 ºC, 500 ºC and 600 ºC using 100 N as the load increment for each step. Discrete steps are seen at RT in Fig. 3(a), where the displacement increases by a small value at each load increment, denoted by the numbers 1, 2, 3, 4 in the plot, keeps constant before the next increment and eventually the coating failed after the final load increment at 400 N (step 4). The step loading curve at 400 ºC shown in Fig. 3(b) is similar to that at RT except that the specimen failed at the fifth load increment. A small displacement rate leading to fracture is found after the fourth load increment at 450 ºC in Fig. 3(c) and the sample failed when the load was increased to 500 N. Fig. 3(d) presents the 8
step loading curve at 500 ºC, small displacement rates are found to exist at the second and third load increment but become largely noticeable at the fourth load increment of 400 N. The displacement rates can be treated as the evidence of creep. This sample was left for creep to fracture at the fourth load increment when an obvious increase in the displacement with time was seen. At 600 ºC, Fig. 3(e), the sample has three load increments and the displacements at each load increment are two or three times larger than the ones at lower temperatures. It can also be observed that a positive displacement rate occurred after the first load increment of 100 N. A large displacement can be seen at the last load increment, which indicates extensive plastic deformation of the material at 600 ºC.
Fig. 3. Small punch multi-step loading curves from RT (21 ºC) to 600 ºC, (a) – (e), at the temperatures indicated in the plots. The numbers in the plots refer to the load step in increments of 100 N.
The instant displacements at each load increment for different temperatures are summarised in Fig. 4. By converting load to stress, displacement to strain and displacement rate to strain rate using Eq. (1) - (3), the corresponding comparison of strain and strain rate in log scale versus the imposed stresses can be obtained as shown in Fig. 5(a) and (b) respectively. It can be seen 9
from Fig. 5(a) specimens exhibit very small and similar strain values in the first three load increments between RT and 500 ºC. But the strain at the second and third load increment at 600 ºC are about one order of magnitude larger than those obtained at lower temperatures, which indicates that the deformation of the alloy increases dramatically. The test at 600 ºC finished after three steps whereas tests at other temperatures ended up with four or five steps. It is found that the strain obtained at 500 ºC and below lie in a good trend with stress increments in Fig. 5(a). The strain rate to stress behaviour below 500 ºC shown in Fig. 5(b) is very close to each other but a separation starts to occur for the test results at 500 ºC and more obvious at 600 ºC. The distinct behaviour in the strain rate at 600 ºC indicates the coating has undergone large plastic deformation.
Fig. 4. Plot of displacement versus load increments of 100 N from RT (21 ºC) to 600 ºC.
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Fig. 5. Plots of equivalent strain versus equivalent stress (a) and minimum strain rate versus equivalent stress (b).
3.3. Fracture surface investigation Fig. 6 shows the macroscopic fracture surfaces after multi-step loading tests at 21 ºC (a), 400 ºC (b), 500 ºC (c) and 600 ºC (d). The fracture surfaces shown at 21 ºC and 400 ºC did not exhibit macroscopic plastic deformation. Fast fracture was found when the final load increment was applied to the two specimens. The fracture surfaces at 500 ºC and 600 ºC show the evidence of cracking which initiated at the centre of the sample and propagated radially outwards to the clamping boundary. Evidence of macroscopic plastic deformation can be seen in Fig. 6(c) and (d).
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Fig. 6. Macroscopic fracture surfaces after SPTs at 21 C (a), 400 C (b) 500 C (c) and 600 C (d).
Fig. 7 depicts the microscopic characteristics of fracture surfaces. Areas of tearing are found from the fracture surface at 21 ºC in Fig. 7(a) due to the inherently ductile nature of γ matrix. Isolated angular β particles, which were proved to be Al rich by chemical analysis, are present on the fracture surface. Flat regions are visible in Fig. 7(a) and these distinctive features are probably the splat boundaries that formed due to the inter-splat shearing during testing. The fracture surface at 400 ºC still exhibits γ matrix “tearing” feature with discrete β particles in Fig. 7(b). Similarly at 500 ºC, γ matrix tearing remains the main deformation mechanism and the evidence of β particles on the fracture surface can also be found. When the temperature increased to 600 ºC, isolated β particles can hardly be observed and extensive tearing has become the dominant failure mode in Fig. 7(d).
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Fig. 7. Microscopic fracture surface characteristics of the CoNiCrAlY coating after multi-step loading SPT at 21 ºC (a), 400 ºC (b), 500 ºC (c) and 600 ºC (d), showing areas of tearing, isolated β particles and the transition of fracture behaviour.
4. Discussion 4.1. Data interpretation For the multi-step loading SPTs shown in Fig. 3, the displacement can be very small when the applied load, i.e. 100 N, is insufficient to produce enough deformation at low temperatures. Thus the percent error can be quite large when calculating the equivalent strain. Fig. 5(a) emphases this and is possibly the reason for the discrepancies of the two low strain points at 60 MPa. Apart from these two points, a general trend can be found, where strain is not very sensitive to temperature but varies with stress according to Eq. (4) when temperature is below 500 °C,
(4)
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It is generally found that the yield strength of MCrAlY coatings is about 500 MPa between 25 – 500 C and decreased to 200 MPa or less above 500 C [46]. Since the stresses applied in SPTs are much smaller than the yield strength at temperatures below 500 °C, the continuous plastic deformation did not occur after each load increment. The instant displacements of the last load increment are similar at 450 C and 500 C, implying that the ductility at both temperatures is close to each other. But it can be seen an increase in the strain rates at 500 °C in Fig. 5(b), which separates from the results at lower temperatures. This separation in the strain rate indicates the ductility of the coating starts to increase and the deformation has become thermally activated at 500 °C. Subsequently the step loading results at 600 C show a much larger increase in displacements and strain, implying that the stress-strain behaviour has become very sensitive to temperature. The strain rate behaviour becomes more distinct at 600 °C compared to that at lower temperatures in Fig. 5(b), exhibiting the further evidence for the ductile-to-brittle transition. It is indicated that the deformation has become plastic flow dominant and undergoes continuous plastic deformation at 600 °C from Fig. 5. According to the work reported by Noebe et al. [24], the yield strength of NiAl is about 250 MPa below 500 C and exhibits a sharp drop between 500 – 600 C. Thus, the stress applied at 600 C in multi-step loading SPTs in Fig. 5 can be very close or even higher to the yield strength. This could produce the extensive increase in displacement and strain, as seen in Fig. 5. Therefore, it can be concluded that the ductile-to-brittle transition of this CoNiCrAlY is likely to occur between 500-600 C, where the coating exhibits the brittle characteristics in deformation below 500 C but becomes prominently plastic deformable at 600 C.
4.2. Fracture mechanisms Since the microstructure of the coating shown in Fig. 2 consists of the γ-phase matrix and monocrystalline β-phase precipitates, it is possible that cracking occurs in the brittle β-phase grains at low temperatures and propagates in the ductile γ-phase matrix. The EBSD phase 14
map in Fig. 2(c) and (d) shows that regions of fine-grained β-phases around the particle boundary which can act as paths for crack initiation and are the reasons for the large flat features observed in Fig. 7(a). One of these paths can be referred to the periphery of the powder particle which has, in part, retained the original powder microstructure. Crack propagation along this path would leave a concave feature in the fracture surface that matches the shape of the powder particle due to the inter-splat shearing, such as the one shown in Fig. 7(a). At low temperatures, the energy required for crack propagation in the β-phase is less than the energy required for plastic flow. Since the volume fraction of the γ phase is about 70%, thus a mixture of γ-matrix tearing and isolated β particles are present, showing the β phases are still brittle. This is the reason why both brittle and ductile behaviour can be observed between 400-500 C in Fig. 7(b) and (c). Since the deformation of the β phase has been thermally activated at 600 C, thus less energy is required for the plastic flow than brittle cleavage in the β-phase. The gradual decrease in the threshold energy for plastic flow allows the β-phase to deform with the γ matrix, as shown in Fig. 7(d), causing the brittle to ductile transition to occur. The fracture surfaces show the tearing is the main failure mode above this DBTT, with a great similarity to the fracture surfaces reported at 750 C [31].
4.3. Comparison with previous studies The CoNiCrAlY coating alloy employed in this study consists of γ matrix phase and β-NiAl secondary phase. The γ matrix phase is ductile at all temperatures and only the β-NiAl phase exhibits the ductile to brittle transition behaviour. The temperature dependent ductility behaviour of β-NiAl reviewed by Noebe et al. [24] shows that the DBTT is between 500 – 700 C. It should be noted that the DBTT of MCrAlY depends on the amount of β-NiAl phase since only β-NiAl exhibits the ductile to brittle transition behaviour [18]. Meetham et al. [3] reported that diffusional aluminides which are predominantly brittle β phase exhibit higher DBTT than those overlay coatings with brittle β and ductile γ phase. Eskner et al. [44] found 15
that the DBTT of a nickel aluminide coating was about 760 C. Since the volume fraction of β-NiAl phase in the CoNiCrAlY coating used in this study is around 30%, the DBTT is expected to be less than 760 C. With regard to the DBTT of MCrAlY alloys, the available data are limited but values in the range 400-900 C are suggested compared to the present range 500-600 C, which shows good agreement among various MCrAlY alloys reviewed by Bose [18]. However, previous studies involve different MCrAlY compositions and phases, thus direct comparison with these studies are not applicable. The loading curves at 400 C and 450 C still exhibits brittle characteristics though the fracture surface indicates some matrix tearing exists due to the presence of ductile matrix. From the evolution of the fracture surfaces in Fig. 7, a clear difference between the fracture behaviour across the test temperature range can be seen, in which γ matrix tearing and isolated β particles were widely observed below 500 C but extensive tearing becomes dominant at 600 °C. Therefore, by combining the test results and fracture characteristics obtained from multi-step loading SPTs, it is believed that the brittle to ductile transition of this CoNiCrAlY coating occurs between 500 ~ 600 °C, where at 500 °C the transition is about to occur and 600 °C and above is clearly ductile tearing dominant.
5. Conclusions
In this work it is shown that the small punch multi-step loading test is a suitable method for qualitatively evaluating the ductile-to-brittle transition behaviour of the CoNiCrAlY bond coat materials.
The displacement and strain obtained at each load increment are small below 500 C, but the resultant strain at 600 C is about one order of magnitude larger. Distinct stress-strain behaviour is observed at 600 C, indicating that the coating has undergone extensive plastic deformation with an obvious increase in ductility.
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Evidence of inter-splat shearing, matrix tearing and isolated β particles were found at the fracture surfaces between RT and 500 C due to the presence of ductile γ-Ni matrix phase and brittle β-NiAl phase. Extensive ductile tearing becomes the dominant failure mode at 600 C due to the brittle to ductile transition occurs between 500 – 600 C.
Acknowledgements The authors would like to thank Mr. Shane Maskill for his skilled assistance in carrying out the small punch testing experiments at the University of Nottingham. The authors also thank Prof. Wei Sun and Prof. Graham McCartney for help discussions. The financial support from Provincial Applied Research Programme for Commonweal Technology funded by Science Technology Department of Zhejiang Province (No. 2016C31023) and Ningbo Natural Science Foundation Programme (No. 2016A610114) are acknowledged.
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