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Microstructure evolution and dislocation behaviour in high chromium, fully ferritic steels strengthened by intermetallic Laves phases
Accepted Manuscript Title: Microstructure Evolution and Dislocation Behaviour in High Chromium, Fully Ferritic Steels Strengthened by Intermetallic Laves Phases Authors: Jennifer Lopez Barrilao, Bernd Kuhn, Egbert Wessel PII: DOI: Reference:
S0968-4328(17)30424-9 https://doi.org/10.1016/j.micron.2018.02.008 JMIC 2533
To appear in:
Micron
Received date: Revised date: Accepted date:
9-11-2017 20-2-2018 20-2-2018
Please cite this article as: Barrilao, Jennifer Lopez, Kuhn, Bernd, Wessel, Egbert, Microstructure Evolution and Dislocation Behaviour in High Chromium, Fully Ferritic Steels Strengthened by Intermetallic Laves Phases.Micron https://doi.org/10.1016/j.micron.2018.02.008 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 proof before it is published in its final 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.
Microstructure Evolution and Dislocation Behaviour in High Chromium, Fully Ferritic Steels Strengthened by Intermetallic Laves Phases Jennifer Lopez Barrilao*, Bernd Kuhn, Egbert Wessel Forschungszentrum Jülich GmbH Institute for Energy and Climate Research Microstructure and Properties of Materials (IEK-2) 52428 Jülich, Germany *email:
[email protected]
Fully ferritic steels as new idea to improve creep resistance. Focus on sub-grain structure, dislocation and particle-free zone evolution. Microstructure and particle evolution in fully ferritic steels. In-depth investigations of dislocation behaviour and its possible effect on the mechanical response.
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Highlights:
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Abstract
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In the present study a stainless, high strength, ferritic (non- martensitic) steel was analysed regarding microstructure and particle evolution. The preceding hot-rolling process of the steel results in the formation of sub-grain structures, which disappear over time at high temperature. Besides that the formation of particle-free zones was observed. The pronounced formation of these zones preferentially appears close to high angle grain boundaries and is considered to be responsible for long-term material failure under creep conditions. The reasons for this are lacking particle hardening and thus a concentration and accumulation of deformation in the particle free areas close to the grain boundaries. Accordingly in-depth investigations were performed by electron microscopy to analyse dislocation behaviour and its possible effect on the mechanical response of these weak areas.
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KEYWORDS: Laves phase, Microstructure, Intermetallic particles, Ferritic steel, Advanced Ultra Supercritical (AUSC) power plants, Dislocations, Particle-free zones
Introduction
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Present investigations focus on a new concept of fully ferritic, stainless steels free of martensitic transformation for the application in high temperature energy conversion systems. The presented trial alloy 2.5W0.57Nb0Ti was produced in the framework of Crofer® 22 H development (commercially available from VDM Metals GmbH, Germany), where more than 50 trial alloys were designed. Crofer® 22 H was designed as an interconnect material for solid oxide fuel cells (SOFCs) for operation temperatures of about 800 °C. In general such steels show promising steam oxidation resistance in the temperature range from 600 °C to 650 °C due to their high chromium content (18 23 wt.%) (Kuhn et al., 2014). In comparison the oxidation resistance of state of the art 9 – 12 % Cr Creep Strength Enhanced Ferritic (CSEF) steels, e.g. P/T91 and P/T92, is limited to temperatures up to
620 °C. Guarantee of sufficient oxidation resistance beyond 620 °C requires higher chromium content (Quadakkers and Zurek, 2010; Wright and Dooley, 2010), but results in promotion of Z-phase formation at the expense of MX (M = Nb, V; X = C, N) particles (Danielsen et al., 2006). Due to the strengthening concept of AFM steels this formation is accompanied by a drop in long-term creep strength, because the strength is mainly obtained by MX (M = Nb, V; X = C, N) and partly by M23C6 (M = Cr) particles as well as solid-solution hardening (Hald, 2008; Maruyama et al., 2001; Viswanathan and Bakker, 2001).
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Experimental
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In contrast to CSEF steels the strength of the presented novel fully ferritic steels is obtained by solidsolution and intermetallic Laves phase particles of the type (Fe,Cr,Si)2(Nb,W), which are controversially discussed in CSEF steels (Hald, 1996; Hosoi et al., 1986). Therefore detrimental interactions of multitudinous precipitate types do not have to be considered. Furthermore the fully ferritic steels possess superior creep behaviour in the temperature range from 600 °C to 650 °C (Kuhn et al., 2014) and therefore may fulfil the current and future requirements, e.g. higher conversion efficiency, lower CO2 emission and higher operational flexibility. Further development and optimisation of trial alloys resulted in so called HiperFer steels (High Performance Ferrite), a new generation of fully ferritic steels with a chromium content of 17 wt.% (Kuhn et al., 2016). Prior investigations of Crofer based alloys were primarily focused on growth or rather evolution of Laves phase particles and the influence of chemical composition on phase formation in the temperature range from 600 °C to 650 °C. By means of high resolution scanning (HR-SEM) and transmission electron microscopy (HR-TEM) it was shown that the strengthening particles are thermodynamically stable over the whole time range covered so far (e. g. approx. 40,000 h at 600 °C). Due to these results an explanation for the differences in creep behaviour of the trial alloys can be given, but does not give explanations about failure mechanisms (Lopez Barrilao, 2017; Lopez Barrilao et al., 2017a; Lopez Barrilao et al., 2017b; Lopez Barrilao et al., 2015). Therefore the most creep resistant trial alloy from Crofer development 2.5W0.57Nb0Ti was chosen to investigate potential creep damage and failure mechanisms, regarding dislocation movement and localization of deformation.
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Chemical Composition and Production of Trial Alloys
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Alloy 2.5W0.57Nb0Ti was produced by vacuum melting of an approximately 10 kg ingot by VDM Metals GmbH and annealed at 1080 °C for 2 h. Afterwards the ingots were processed to hot-rolled plates at approx. 980 °C and cooled to ambient temperature. The chemical composition of alloy 2.5W0.57Nb0Ti is given in Table 1. Table 1: Chemical composition of the trial alloy 2.5W0.57Nb0Ti (in wt.%).
2.5W0.57Nb0Ti
C
N
Cr
W
Nb
Si
Mn
La
Ti
0.004
0.008
22.95
2.50
0.57
0.20
0.46
0.034
0.004
Sample Annealing and Preparation
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Cubic samples (approx. 5 mm edge length, cut from the hot-rolled plates) were annealed in a preheated furnace at 650 °C for 2 h, 100 h and 1,000 h, followed by water quenching. Long-term aging samples were taken from creep specimens perpendicular to the loading direction. The samples entered the creep test in the hot-rolled state without additional heat treatment before testing (rupture times: 2,425 h at 100 MPa and 9,808 h at 70 MPa). Observations of particle size evolution and dislocation behaviour without the influence of mechanical stress were performed at samples extracted from the quasi stress-free head sections of the creep specimens (cf. Figure 1, position 1). For consideration of the impact of creep stress on dislocation behaviour samples were taken from the uniformly deformed gauge length (cf. Figure 1, position 2)
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Figure 1: Creep specimen: Before testing (bottom) and after rupture (top). The positions where the samples were cut from are indicate by rectangles: position 1 represents the stress-free head section, position 2 the uniformly deformed gauge length section (Lopez Barrilao, 2017).
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For the investigations the samples were metallographically prepared. The preparation consisted of mounting in epoxy resin, grinding and polishing to sub-micron finish for 1 - 2 h per polishing step. Samples for particle and microstructure analyses were additionally etched at 1.5 V in 5% H2SO4 to enhance the particles/matrix contrast.
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Dislocations were analysed using a Zeiss Libra 200 transmission electron microscope (TEM) with an acceleration voltage of 200 kV. All images were taken in scanning TEM (STEM) mode applying a high angle annular dark field (HAADF) detector. Energy dispersive X-ray (EDX) spectroscopy (X-Max 80, Oxford Instruments) was used for the chemical composition analyses.
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The thickness of the investigated lamellae were approx. 150 nm with a size of approx. 10 µm × 10µm. Preparation of the lamellae was performed by a Zeiss Auriga cross beam focused ion beam (FIB) with Ga-ion beam. HR-FESEM High resolution field emission scanning electron microscopy (HR-FESEM) images were taken by a Zeiss Merlin SEM (2.5k- and 10k-fold magnification, 6144 × 4608 pixels, 7.44 nm per pixel in 2.5k / 1.86 nm per pixel in 10k magnification, back scattered electrons) for the evaluation of particle size distribution and microstructure evolution. The images were analysed by applying the commercial
software package AnalysisPro®. Electron backscatter diffraction (EBSD) has been used to analyse grain boundaries and deformations.
Results and Discussions Particle Size Evolution at 650 °C
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The investigation of particle size evolution is not the main focus of this study and will therefore be reiterated shortly only. For detailed analysis of several trial alloys refer to (Lopez Barrilao, 2017; Lopez Barrilao et al., 2017a; Lopez Barrilao et al., 2017b) . The evolution of the equivalent circle diameter (ECD) distribution of the strengthening particles in the creep strongest alloy 2.5W0.57Nb0Ti at 650 °C is given in Figure 2. The results were evaluated from the images in Figure 3 or rather similar images of other sample locations. Figure 2 displays the size classification of Laves phase precipitates in alloy 2.5W0.57Nb0Ti with regard to particle area fraction (fraction of particles in the respective size class proportional to the overall area covered by all particles). Former studies of particles size development in the 2.5W0.57Nb0Ti material indicated a tendency to maintain smaller particles than other trial alloys from Crofer development. This finding was ascribed to improved particle stability and/or dynamic precipitation processes, which in turn was supposed to be the main reason for the superior creep performance (cf. (Kuhn et al., 2014)) of this trial steel. However, the particle evolution investigations did not provide satisfactory information concerning potential creep damage and failure mechanisms.
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Figure 2: Development of particle size distributions in terms of equivalent circle diameter (ECD) (Lopez Barrilao et al., 2016).
Sub-grain Structure Evolution After different annealing time steps the observations show evolution of two types of sub-grain structures besides expected particle coarsening. Figure 3 illustrates the microstructure evolution after different annealing steps from the as received condition up to 9,808 h of annealing at 650 °C. The apparent sub-grain structures are a result of the preceding hot-rolling process and are typical for fully ferritic, Laves phase strengthened steels, which were thermomechanically processed (Talik, 2016). Figure 3a displays the as received state (as rolled, 0 h of annealing), where sub-grain
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structures are visible across the whole microstructure. These initial sub-grain boundaries as well as the high-angle grain boundaries are occasionally decorated by particles. Both boundary types were progressively occupied by precipitates up to 1,000 h of annealing (cf. Figure 3b-d). Coarsening of the boundary particles continues in case of high-angle grain boundaries only (after 2,425 h and 9,808 h in the head section (HS) and gauge length (GL) section, cf. Figure 3e-h), while the sub-grain structures start to disappear between 1,000 h and 2,425 h (compare Figure 3d to Figure 3e-h). The initial subgrain structures are rarely retained after 2,425 h and appear to be completely absent after 9,808 h of aging. Comparisons of the HS images after 2,425 h and 9,808 h (cf Figure 3e+g) with the corresponding images of the gauge length sections (cf. Figure 3f+h) do not show substantial differences regarding evolution of the initial the sub-grain structures. Therefore only temperature and time or rather thermal recovery are considered to contribute to disintegration of the initial subgrain structures. Consequently the influence of mechanical stress is considered to be a minor one.
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Figure 3: HR-FESEM images after different annealing steps at 650 °C: a) 0 h (as received state after hot rolling) (Ning, 2013), b) 2 h, c) 100 h, d) 1,000 h, e) 2,425 h (head section), f) 2,425 h (gauge length), g) 9,808 h (head section) and h) 9,808 h (gauge length)(Lopez Barrilao et al., 2016).
Evolution of Particle-Free Zones Apart from the evolution and final disappearance of the initial sub-grain structures, the formation of particle-free zones (PFZ) can be observed (cf. Figure 3). These zones preferentially form along the high-angle grain boundaries and thus are potentially linked to void formation and thus to creep damage and failure (cf. Figure 4). Because of that these zones are of particular interest.
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Figure 4: Void formation during creep testing at 650 °C in the gauge lengths of alloy 2.5W0.57Nb0Ti: a) after 2,524 h at 100 MPa and b) after 9,808 h at 70 MPa.
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A comparison of the images after different annealing times in Figure 3 shows an increase of their width over time. The reason for this is Ostwald ripening, where particles on high-angle grain boundaries coarsen on the expense of small matrix particles close to these boundaries (Gottstein, 2007). Mean width development and standard deviation of the particle-free zones are given in Figure 5. The calculation of the mean and standard deviation were based on measurements of the furnace annealed samples (up to 1,000 h) and samples after creep testing (2,425 h at 100 MPa and 9,808 h at 70 MPa) at a minimum of 11 different locations along the high-angle grain boundaries.
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Figure 5: Mean width and standard deviation of particle-free zones at high angle grain boundaries after different annealing steps at 650 °C.
The measurements of the PFZ width confirm a growth over time. Up to 2,425 h the width of the zones increases progressively without strong variations between the different locations of the measurements (0.45 ± 0.11 µm after 2 h up to 1.72 ± 0.38 µm after 2,425 h; HS). Besides this increase of the width, the measurement show variations between the head and gauge length sections (1.72 ± 0.38 µm; HS and 1.58 ± 0.35 µm; GL after 2,425 h of annealing or rather 4.25 ± 0.63 µm; HS and 2.74 ± 1.00 µm; GL after 9,808 h of annealing). Furthermore the width of gauge length sections vary stronger after longer annealing (4.25 ± 0.63 µm; HS and 2.74 ± 1.00 µm; GL after 9,808 h). Smaller zone width in case of gauge length sections in comparison to the head sections may be
caused by strain localisation and thus increased dislocation densities. This consequently potentially boosts heterogeneous precipitation due to lower nucleation energy in these areas if an oversaturated matrix is provided (1984). Detailed analyses of dislocation behaviour in the particle free zones will be discussed hereafter.
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As mentioned before the Crofer® 22 H material is strengthened by solid-solution and intermetallic Laves phase particles. The predecessor of this material is the commercially available material Crofer® 22 APU. This steel type is strengthened by solid solution only. Thus, the increased strength of the Crofer® 22 H material is mainly the result of additional particle hardening (also see (Froitzheim et al., 2008; Kuhn et al., 2011)). Therefore the observed particle-free zones are of specific interest, because in these areas the 2.5W0.57Nb0Ti material may rather behave like a solely solid-solution strengthened material. In-depth analyses of the PFZs were performed by HR-FESEM using EBSD technique for crystal (mis-)orientation mappings and HR-TEM for investigations of dislocation behaviour and dislocation densities. Band contrast images and corresponding misorientation mappings of alloy 2.5W0.57Nb0Ti in the as received condition (0 h of annealing, cf. Figure 6a+b), after 2,425 h (cf. Figure 6c-f) and after 9,808 h (cf. Figure 6g-j) of aging from the head and gauge length sections are given in Figure 6.
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Figure 6: Alloy 2.5W0.57Nb0Ti at 650 °C: a) EBSD band contrast image with grain boundaries of the as received condition (black: high-angle, red: low angle), b) misorientation mapping of the as received condition, c) EBSD band contrast image with grain boundaries image after 2,425 h of aging (head section, black: high-angle, red: low angle), d) misorientation mapping after 2,425 h (head section), e) EBSD band contrast image with grain boundaries after 2,425 h (gauge length, black: high-angle, red: low angle), f) misorientation mapping after 2,425 h (gauge length), g) grain EBSD band contrast image with grain boundaries after 9,808 h (head section, black: high-angle, red: low angle), d) misorientation mapping after 9,808 h (head section), e) EBSD band contrast image with grain boundaries after 9,808 h (gauge length, black: high-angle, red: low angle) and f) misorientation mapping after 9,808 h (gauge length) (Lopez Barrilao, 2017).
The band contrast images of the quasi stress-free head section and the stressed gauge length section show pronounced formation of low-angle grain boundaries within the PFZs. These newly formed low-
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angle grain boundaries interestingly are not occupied by particles. The reason for the formation of these boundaries is lattice deformation caused by external and/or residual stress or rather accumulation of dislocations. These low-angle boundaries are only verifiable by EBSD measurements. In general the formed boundaries in the gauge length samples show predominant orientation perpendicular to the high-angle boundaries, which are indicated in black in the left image column of Figure 6. By means of local misorientation mappings (cf. right column of Figure 6) even small differences in crystallographic lattice orientation can be displayed. These small changes in orientation are visible as misorientations of around 1° indicated as green lines. Most of the observed areas (the bcc matrix) certainly do not show strong orientation differences. Therefore most of the area is indicated in blue or rather as nearly 0° in misorientation. In comparison to the PFZ areas, greater sections of misorientation were found in the areas occupied by particles. On the contrary the misorientations in the PFZs are only marginal. Misorientations of around 1° are already visible in the as received state (cf. Figure 6a+b) without additional influence of particles, which are caused by the preceding hot-rolling process. A direct comparison of the mappings in the right column of Figure 6 indicates more pronounced crystal lattice relaxation within the PFZs. This relaxation process is only reduced due to particles. Therefore stronger misorientations arise by particles and thus affect deformation of the lattice, which might be enhanced by the presence of dislocations. These dislocations pin at particles and trigger stronger misorientations in these areas. Additional investigations of the PFZs were performed by TEM. The results provide further detail information on dislocation behaviour in these zones. Figure 7 shows an overview of TEM lamellae, which were cut from the stress-free head section and the stressed gauge length after 2,425 h of creep testing at 650 °C / 100 MPa and after 9,808 h at 70 MPa. The white lines in the left column indicate the borders of the particle-free zones. Magnifications of the areas marked by red squares are given in the right column of Figure 7. The areas indicated by dashed lines in the PFZs in the right column represent the areas where dislocation densities were calculated. The corresponding dislocation densities are given in Table 2. Density calculations only consider freely moveable and solitary dislocations. Dislocations which are already arranged in dislocation structures were neglected.
2,425 h (GL) ≈34.72
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Table 2: Dislocation densities (in 1012 m-2) calculated from the dashed marked areas in Figure 7 of the stress-free head sections and the stressed gauge lengths after 2,425 h at 100 MPa and 9,808 h at 70 MPa.
9,808 h (HS) ≈10.66
9,808 h (GL) ≈53.00
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Comparison of dislocation density calculations of the head section samples after 2,425 h and 9,808 h shows similar values, which indicates a stable state after relatively short time in the absence of stress. The influence of stress results in an increase of dislocation density. An increase in direct comparison between head section and gauge length samples was determined. After 9,808 h of creep testing at 70 MPa the density increases by approx. five times and after 2,425 h at 100 MPa an increase by a factor of three was identified. Thus, the dislocation densities of the gauge length sections apparently show an inverse dependency on the stress level, but a direct proportionality to the particle-free zones width. The dependencies of dislocation densities are given in Figure 8.
IP T SC R U N A M ED PT CC E A Figure 7: STEM (at 200 kV) overview bright field images (left column) and magnifications of the areas marked by red squares (right column) of alloy 2.5W0.57Nb0Ti after 2,425 h a)+b) head section, c)+d) gauge length and after 9,808 h at 650 °C e)+f) head section and g)+h) gauge length. Dislocation densities were calculated from the areas marked by dashed lines (right column). White lines indicate the borders of PFZs (left column) (Lopez Barrilao, 2017; Lopez Barrilao et al., 2016).
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Figure 8: Dependencies of dislocation density on particle-free zone width and stress level.
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In case of the stressed gauge length sections the examination of dislocations by TEM clearly shows a preferable orientation of the dislocations parallel to the grain boundaries (cf. Figure 7c+d and g+h). This orientation may support an accumulation of dislocations at the grain boundaries and consequently benefit void formation (Gottstein, 2007) and finally premature failure. Furthermore the preferable orientation decreases diffusion towards grain boundaries, but promotes diffusion alongside the dislocations and thus may decelerate the growth of grain boundary particles. EDX measurements of the matrix within the PFZ exhibits a tungsten content in the gauge length of approx. 1.9 wt.% in case of the 2,425 h sample tested at 100 MPa and approx. 0.65 wt.% in case of the 70 MPa sample after 9,808 h. Consequently only a third of the tungsten content remains after a four times longer testing period in the PFZs of the 70 MPa sample with enhanced dislocation density (cf. Table 2), while parts of the tungsten content conduce constantly to e.g. particle coarsening (cf. Figure 2). The tungsten contents in the PFZs in the quasi stress-free head sections do not vary significantly from the ones found in the gauge length. The head section of the 2,425 h sample tested at 100 MPa contains tungsten contents between 1.9 wt.% and 2.1 wt.%, while in case of the 9,808 h sample tested at 70 MPa tungsten contents between 0.8 wt.% and 1.0 wt.% were found. Therefor the reduction of the tungsten content is mainly a consequence of longer testing duration. The reduction in tungsten content and thus reduced solid-solution strengthening of the PFZs however might give an explanation for localized, increased deformation in the PFZs encountered in the 70 MPa sample tested for 9,808 h. In addition this influences the nucleation of Laves phase particles, which is strongly increased by enhanced dislocation densities (Hsiao et al.; Kuhn et al., 2016). Therefore the reason for smaller PFZ widths in case of gauge length sections in comparison to head sections is proposed to be a combination of preferable dislocation orientation and thus decelerated diffusion towards grain boundaries and the accompanied influence on particle nucleation and growth (cf Table 3). Final conclusions on the relations between particle-free zones, their width, dislocation densities and damage formation or rather premature failure cannot be drawn yet. More experiments are needed in order to obtain enough statistical coverage. Ongoing research focuses on the provision of satisfactory data as well as chemical composition and heat treatment optimisations to minimise the formation of PFZs.
Table 3: Mean width development and standard deviation of particle-free zones after different annealing times at 650 °C (in µm).
2h 0.45±0.11
100 h 0.63±0.18
1,000 h 0.87±0.25
2,425 h (HS) 1.72±0.38
2,425 h (GL) 1.58±0.35
9,808 h (HS) 4.25±0.63
9,808 h (GL) 2.74±0.1.00
Summary and Conclusion
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Preceding investigation of high chromium, fully ferritic steels, strengthened by intermetallic Laves phase particles contextualise particle evolution and creep behaviour (Lopez Barrilao, 2017; Lopez Barrilao et al., 2017a; Lopez Barrilao et al., 2017b; Lopez Barrilao et al., 2015). In comparison to all analysed trial alloys, the creep strongest alloy 2.5W0.57Nb0Ti (Kuhn et al., 2014) in general exhibits the tendency to maintain smaller particles during long-term high temperature application. This result was attributed to improved particle stability and/or dynamic precipitation processes causing superior creep performance. But these particle investigations did not give information about the reasons for material damage and failure. Therefore alloy 2.5W0.57Nb0Ti was analysed concerning sub-grain structure and dislocation evolution. Initial sub-grain boundaries become decorated by Laves phase particles during short-term annealing, but continually decompose until disappearance by recovery during mid-term annealing. These initial sub-grain structures are a result of the preceding hot-rolling process, with precipitates gradually occupying the cell boundaries during aging and are not to be confused with low-angle boundaries developing in particle free zones along high-angle grain boundaries during creep deformation. Comparisons of samples taken from quasi stress-free head sections and stressed gauge length sections after creep testing do not show significant differences regarding the evolution of the initial sub-grain structures. Therefore additional influence of stress on these can be assessed to be as a minor one.
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Additionally the formation of particle-free zones along high-angle grain boundaries was observed. These zones are of specific interest, because within these the material acts similar to a solely solidsolution strengthened material, what in comparison to the particle populated grain interiors is accompanied by a drop in creep strength in the long-term. The analyses show that the particle-free zones widths increase with time, but decrease with stress. Furthermore transmission electron microscopy investigations were performed to obtain detailed information about the PFZs. Increased dislocation densities as well as preferred orientation of the dislocations parallel to high-angle grain boundaries were verified. This accumulation of dislocations may benefit creep void formation and thus premature material failure. Moreover their preferred orientation may decelerate diffusion towards the grain boundaries and thus influence boundary particle growth. Additional chemical analyses show strong variations of the tungsten content depending on dislocation densities. Therefore changes in solid-solution strengthening in the PFZs can be assumed. Due to these changes in solid-solution strengthening, changes in dislocation densities and diffusion rates the particle-free zones are suspected to be a main factor contributing to creep damage and failure.
Acknowledgments The authors would like to thank B. Werner, H. Reiners, D. Liebert, A. Moser, M. Braun and W. Lange for mechanical (long-term) testing and sample annealing. Moreover the authors gratefully acknowledge the support of D. Eßer, V. Gutzeit and J. Bartsch for sample and metallographic
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preparation as well as J. Ning and W. Chen for the provided images. The support of VDM Metals GmbH with experimental material is greatly acknowledged.
References
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1984. Werkstoffkunde Stahl, Band 1: Grundlagen. Verein Deutscher Eisenhüttenleute, SpringerVerlag Berlin Heidelberg Danielsen, H.K., Hald, J., Grumsen, F.B., Somers, M.A.J., 2006. On the crystal structure of Z-phase Cr (V, Nb) N. Metallurgical and Materials Transactions A 37, 2633-2640. Froitzheim, J., Meier, G.H., Niewolak, L., Ennis, P.J., Hattendorf, H., Singheiser, L., Quadakkers, W.J., 2008. Development of high strength ferritic steel for interconnect application in SOFCs. Journal of Power Sources 178, 163-173. Gottstein, G., 2007. Physikalische Grundlagen der Materialkunde. Springer Verlag. Hald, J., 1996. Metallurgy and creep properties of new 9-12%Cr steels. Steel Research 67, 369-374. Hald, J., 2008. Microstructure and long-term creep properties of 9–12% Cr steels. International Journal of Pressure Vessels and Piping 85, 30-37. Hosoi, Y., Wade, N., Kunimitsu, S., Urita, T., 1986. Precipitation behaviour of Laves phase and it's effect on toughness of 10Cr–2Mo ferritic-martensitic steel. Journal of Nuclear Materials 141-143, 461-467. Hsiao, Z.W., Kuhn, B., Yang, S.M., Yang, L.C., Huang, S.Y., Singheiser, L., Kuo, J.C., Lin, D.Y.L.D.Y., The influence of deformation on the precipitation behavior of a ferritic stainless steel, in: LecomteBeckers, J., Dedry, O., Oakey, J., Kuhn, B. (Eds.), 10th Liege Conference : Materials for Advanced Power Engineering 2014. Kuhn, B., Asensio Jimenez, C., Niewolak, L., Hüttel, T., Beck, T., Hattendorf, H., Singheiser, L., Quadakkers, W.J., 2011. Effect of Laves phase strengthening on the mechanical properties of high Cr ferritic steels for solid oxide fuel cell interconnect application. Materials Science and Engineering: A 528, 5888-5899. Kuhn, B., Talík, M., Lopez Barrilao, J., Singheiser, L., Yamamoto, Y., Schmitz-Niederau, M., 2016. Development Status of High performance Ferritic (HiperFer) steels, 8th International Conference on Advances in Materials Technology for Fossil Power Plants, pp. 1019-1028. Kuhn, B., Talík, M., Niewolak, L., Zurek, J., Hattendorf, H., Ennis, P.J., Quadakkers, W.J., Beck, T., Singheiser, L., 2014. Development of high chromium ferritic steels strengthened by intermetallic phases. Materials Science and Engineering A, 594, 372-380. Lopez Barrilao, J., 2017. Microstructure Evolution of Laves Phase Strengthened Ferritic Steels for High Temperature Applications, Schriften des Forschungszentrums Jülich Reihe Energie & Umwelt / Energy & Environment. RWTH Aachen University, Jülich, pp. XVI, 134 pp. Lopez Barrilao, J., Kuhn, B., Wessel, E., 2016. Microstructure and Intermetallic Particle Evolution in Fully Ferritic Steels, 8th International Conference on Advances in Materials Technology for Fossil Power Plants, pp. 1029-1037. Lopez Barrilao, J., Kuhn, B., Wessel, E., 2017a. Identification, size classification and evolution of Laves phase precipitates in high chromium, fully ferritic steels. Micron 101, 221-231. Lopez Barrilao, J., Kuhn, B., Wessel, E., Talík, M., 2017b. Microstructure of intermetallic particle strengthened high-chromium fully ferritic steels. Materials science and technology 33, 1056-1064. Lopez Barrilao, J., Talík, M., Kuhn, B., Wessel, E., 2015. Microstructure Evolution of High Temperature Fully Ferritic Steels, 1st International Conference on Advanced High-Temperature Materials Technology for Sustainable and Reliable Power Engineering (123 HiMAT-2015), Sapporo, Japan. 123rd Committee of Japan Socity for the Promotion of Science, IAI, pp. 15-18. Maruyama, K., Sawada, K., Koike, J., 2001. Strengthening mechanisms of creep resistant tempered martensitic steel. ISIJ international 41, 641-653. Ning, J., 2013. Influence of chemical composition on precipitation reactions in high chromium ferritic steels. RWTH Aachen – Institute of Ferrous Metallurgy. Quadakkers, W.J., Zurek, J., 2010. Oxidation in steam and steam/hydrogen environments. Shreir's Corrosion 1, 407-456.
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Talik, M., 2016. Influence of Initial Thermomechanical Treatment on High Temperature Properties of Laves Phase Strengthened Ferritic Steels, Schriften des Forschungszentrums Jülich Reihe Energie & Umwelt / Energy & Environment. RWTH Aachen University, Jülich, pp. xxiii, 130. Viswanathan, R., Bakker, W., 2001. Materials for ultrasupercritical coal power plants—Boiler materials: Part 1. Journal of Materials Engineering and Performance 10, 81-95. Wright, I.G., Dooley, R.B., 2010. A review of the oxidation behaviour of structural alloys in steam. International Materials Reviews 55, 129-167.
S0968-4328(17)30424-9 https://doi.org/10.1016/j.micron.2018.02.008 JMIC 2533
To appear in:
Micron
Received date: Revised date: Accepted date:
9-11-2017 20-2-2018 20-2-2018
Please cite this article as: Barrilao, Jennifer Lopez, Kuhn, Bernd, Wessel, Egbert, Microstructure Evolution and Dislocation Behaviour in High Chromium, Fully Ferritic Steels Strengthened by Intermetallic Laves Phases.Micron https://doi.org/10.1016/j.micron.2018.02.008 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 proof before it is published in its final 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.
Microstructure Evolution and Dislocation Behaviour in High Chromium, Fully Ferritic Steels Strengthened by Intermetallic Laves Phases Jennifer Lopez Barrilao*, Bernd Kuhn, Egbert Wessel Forschungszentrum Jülich GmbH Institute for Energy and Climate Research Microstructure and Properties of Materials (IEK-2) 52428 Jülich, Germany *email:
[email protected]
Fully ferritic steels as new idea to improve creep resistance. Focus on sub-grain structure, dislocation and particle-free zone evolution. Microstructure and particle evolution in fully ferritic steels. In-depth investigations of dislocation behaviour and its possible effect on the mechanical response.
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Highlights:
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Abstract
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In the present study a stainless, high strength, ferritic (non- martensitic) steel was analysed regarding microstructure and particle evolution. The preceding hot-rolling process of the steel results in the formation of sub-grain structures, which disappear over time at high temperature. Besides that the formation of particle-free zones was observed. The pronounced formation of these zones preferentially appears close to high angle grain boundaries and is considered to be responsible for long-term material failure under creep conditions. The reasons for this are lacking particle hardening and thus a concentration and accumulation of deformation in the particle free areas close to the grain boundaries. Accordingly in-depth investigations were performed by electron microscopy to analyse dislocation behaviour and its possible effect on the mechanical response of these weak areas.
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KEYWORDS: Laves phase, Microstructure, Intermetallic particles, Ferritic steel, Advanced Ultra Supercritical (AUSC) power plants, Dislocations, Particle-free zones
Introduction
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Present investigations focus on a new concept of fully ferritic, stainless steels free of martensitic transformation for the application in high temperature energy conversion systems. The presented trial alloy 2.5W0.57Nb0Ti was produced in the framework of Crofer® 22 H development (commercially available from VDM Metals GmbH, Germany), where more than 50 trial alloys were designed. Crofer® 22 H was designed as an interconnect material for solid oxide fuel cells (SOFCs) for operation temperatures of about 800 °C. In general such steels show promising steam oxidation resistance in the temperature range from 600 °C to 650 °C due to their high chromium content (18 23 wt.%) (Kuhn et al., 2014). In comparison the oxidation resistance of state of the art 9 – 12 % Cr Creep Strength Enhanced Ferritic (CSEF) steels, e.g. P/T91 and P/T92, is limited to temperatures up to
620 °C. Guarantee of sufficient oxidation resistance beyond 620 °C requires higher chromium content (Quadakkers and Zurek, 2010; Wright and Dooley, 2010), but results in promotion of Z-phase formation at the expense of MX (M = Nb, V; X = C, N) particles (Danielsen et al., 2006). Due to the strengthening concept of AFM steels this formation is accompanied by a drop in long-term creep strength, because the strength is mainly obtained by MX (M = Nb, V; X = C, N) and partly by M23C6 (M = Cr) particles as well as solid-solution hardening (Hald, 2008; Maruyama et al., 2001; Viswanathan and Bakker, 2001).
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Experimental
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In contrast to CSEF steels the strength of the presented novel fully ferritic steels is obtained by solidsolution and intermetallic Laves phase particles of the type (Fe,Cr,Si)2(Nb,W), which are controversially discussed in CSEF steels (Hald, 1996; Hosoi et al., 1986). Therefore detrimental interactions of multitudinous precipitate types do not have to be considered. Furthermore the fully ferritic steels possess superior creep behaviour in the temperature range from 600 °C to 650 °C (Kuhn et al., 2014) and therefore may fulfil the current and future requirements, e.g. higher conversion efficiency, lower CO2 emission and higher operational flexibility. Further development and optimisation of trial alloys resulted in so called HiperFer steels (High Performance Ferrite), a new generation of fully ferritic steels with a chromium content of 17 wt.% (Kuhn et al., 2016). Prior investigations of Crofer based alloys were primarily focused on growth or rather evolution of Laves phase particles and the influence of chemical composition on phase formation in the temperature range from 600 °C to 650 °C. By means of high resolution scanning (HR-SEM) and transmission electron microscopy (HR-TEM) it was shown that the strengthening particles are thermodynamically stable over the whole time range covered so far (e. g. approx. 40,000 h at 600 °C). Due to these results an explanation for the differences in creep behaviour of the trial alloys can be given, but does not give explanations about failure mechanisms (Lopez Barrilao, 2017; Lopez Barrilao et al., 2017a; Lopez Barrilao et al., 2017b; Lopez Barrilao et al., 2015). Therefore the most creep resistant trial alloy from Crofer development 2.5W0.57Nb0Ti was chosen to investigate potential creep damage and failure mechanisms, regarding dislocation movement and localization of deformation.
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Chemical Composition and Production of Trial Alloys
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Alloy 2.5W0.57Nb0Ti was produced by vacuum melting of an approximately 10 kg ingot by VDM Metals GmbH and annealed at 1080 °C for 2 h. Afterwards the ingots were processed to hot-rolled plates at approx. 980 °C and cooled to ambient temperature. The chemical composition of alloy 2.5W0.57Nb0Ti is given in Table 1. Table 1: Chemical composition of the trial alloy 2.5W0.57Nb0Ti (in wt.%).
2.5W0.57Nb0Ti
C
N
Cr
W
Nb
Si
Mn
La
Ti
0.004
0.008
22.95
2.50
0.57
0.20
0.46
0.034
0.004
Sample Annealing and Preparation
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Cubic samples (approx. 5 mm edge length, cut from the hot-rolled plates) were annealed in a preheated furnace at 650 °C for 2 h, 100 h and 1,000 h, followed by water quenching. Long-term aging samples were taken from creep specimens perpendicular to the loading direction. The samples entered the creep test in the hot-rolled state without additional heat treatment before testing (rupture times: 2,425 h at 100 MPa and 9,808 h at 70 MPa). Observations of particle size evolution and dislocation behaviour without the influence of mechanical stress were performed at samples extracted from the quasi stress-free head sections of the creep specimens (cf. Figure 1, position 1). For consideration of the impact of creep stress on dislocation behaviour samples were taken from the uniformly deformed gauge length (cf. Figure 1, position 2)
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Figure 1: Creep specimen: Before testing (bottom) and after rupture (top). The positions where the samples were cut from are indicate by rectangles: position 1 represents the stress-free head section, position 2 the uniformly deformed gauge length section (Lopez Barrilao, 2017).
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For the investigations the samples were metallographically prepared. The preparation consisted of mounting in epoxy resin, grinding and polishing to sub-micron finish for 1 - 2 h per polishing step. Samples for particle and microstructure analyses were additionally etched at 1.5 V in 5% H2SO4 to enhance the particles/matrix contrast.
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Dislocations were analysed using a Zeiss Libra 200 transmission electron microscope (TEM) with an acceleration voltage of 200 kV. All images were taken in scanning TEM (STEM) mode applying a high angle annular dark field (HAADF) detector. Energy dispersive X-ray (EDX) spectroscopy (X-Max 80, Oxford Instruments) was used for the chemical composition analyses.
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The thickness of the investigated lamellae were approx. 150 nm with a size of approx. 10 µm × 10µm. Preparation of the lamellae was performed by a Zeiss Auriga cross beam focused ion beam (FIB) with Ga-ion beam. HR-FESEM High resolution field emission scanning electron microscopy (HR-FESEM) images were taken by a Zeiss Merlin SEM (2.5k- and 10k-fold magnification, 6144 × 4608 pixels, 7.44 nm per pixel in 2.5k / 1.86 nm per pixel in 10k magnification, back scattered electrons) for the evaluation of particle size distribution and microstructure evolution. The images were analysed by applying the commercial
software package AnalysisPro®. Electron backscatter diffraction (EBSD) has been used to analyse grain boundaries and deformations.
Results and Discussions Particle Size Evolution at 650 °C
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The investigation of particle size evolution is not the main focus of this study and will therefore be reiterated shortly only. For detailed analysis of several trial alloys refer to (Lopez Barrilao, 2017; Lopez Barrilao et al., 2017a; Lopez Barrilao et al., 2017b) . The evolution of the equivalent circle diameter (ECD) distribution of the strengthening particles in the creep strongest alloy 2.5W0.57Nb0Ti at 650 °C is given in Figure 2. The results were evaluated from the images in Figure 3 or rather similar images of other sample locations. Figure 2 displays the size classification of Laves phase precipitates in alloy 2.5W0.57Nb0Ti with regard to particle area fraction (fraction of particles in the respective size class proportional to the overall area covered by all particles). Former studies of particles size development in the 2.5W0.57Nb0Ti material indicated a tendency to maintain smaller particles than other trial alloys from Crofer development. This finding was ascribed to improved particle stability and/or dynamic precipitation processes, which in turn was supposed to be the main reason for the superior creep performance (cf. (Kuhn et al., 2014)) of this trial steel. However, the particle evolution investigations did not provide satisfactory information concerning potential creep damage and failure mechanisms.
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Figure 2: Development of particle size distributions in terms of equivalent circle diameter (ECD) (Lopez Barrilao et al., 2016).
Sub-grain Structure Evolution After different annealing time steps the observations show evolution of two types of sub-grain structures besides expected particle coarsening. Figure 3 illustrates the microstructure evolution after different annealing steps from the as received condition up to 9,808 h of annealing at 650 °C. The apparent sub-grain structures are a result of the preceding hot-rolling process and are typical for fully ferritic, Laves phase strengthened steels, which were thermomechanically processed (Talik, 2016). Figure 3a displays the as received state (as rolled, 0 h of annealing), where sub-grain
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structures are visible across the whole microstructure. These initial sub-grain boundaries as well as the high-angle grain boundaries are occasionally decorated by particles. Both boundary types were progressively occupied by precipitates up to 1,000 h of annealing (cf. Figure 3b-d). Coarsening of the boundary particles continues in case of high-angle grain boundaries only (after 2,425 h and 9,808 h in the head section (HS) and gauge length (GL) section, cf. Figure 3e-h), while the sub-grain structures start to disappear between 1,000 h and 2,425 h (compare Figure 3d to Figure 3e-h). The initial subgrain structures are rarely retained after 2,425 h and appear to be completely absent after 9,808 h of aging. Comparisons of the HS images after 2,425 h and 9,808 h (cf Figure 3e+g) with the corresponding images of the gauge length sections (cf. Figure 3f+h) do not show substantial differences regarding evolution of the initial the sub-grain structures. Therefore only temperature and time or rather thermal recovery are considered to contribute to disintegration of the initial subgrain structures. Consequently the influence of mechanical stress is considered to be a minor one.
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Figure 3: HR-FESEM images after different annealing steps at 650 °C: a) 0 h (as received state after hot rolling) (Ning, 2013), b) 2 h, c) 100 h, d) 1,000 h, e) 2,425 h (head section), f) 2,425 h (gauge length), g) 9,808 h (head section) and h) 9,808 h (gauge length)(Lopez Barrilao et al., 2016).
Evolution of Particle-Free Zones Apart from the evolution and final disappearance of the initial sub-grain structures, the formation of particle-free zones (PFZ) can be observed (cf. Figure 3). These zones preferentially form along the high-angle grain boundaries and thus are potentially linked to void formation and thus to creep damage and failure (cf. Figure 4). Because of that these zones are of particular interest.
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Figure 4: Void formation during creep testing at 650 °C in the gauge lengths of alloy 2.5W0.57Nb0Ti: a) after 2,524 h at 100 MPa and b) after 9,808 h at 70 MPa.
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A comparison of the images after different annealing times in Figure 3 shows an increase of their width over time. The reason for this is Ostwald ripening, where particles on high-angle grain boundaries coarsen on the expense of small matrix particles close to these boundaries (Gottstein, 2007). Mean width development and standard deviation of the particle-free zones are given in Figure 5. The calculation of the mean and standard deviation were based on measurements of the furnace annealed samples (up to 1,000 h) and samples after creep testing (2,425 h at 100 MPa and 9,808 h at 70 MPa) at a minimum of 11 different locations along the high-angle grain boundaries.
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Figure 5: Mean width and standard deviation of particle-free zones at high angle grain boundaries after different annealing steps at 650 °C.
The measurements of the PFZ width confirm a growth over time. Up to 2,425 h the width of the zones increases progressively without strong variations between the different locations of the measurements (0.45 ± 0.11 µm after 2 h up to 1.72 ± 0.38 µm after 2,425 h; HS). Besides this increase of the width, the measurement show variations between the head and gauge length sections (1.72 ± 0.38 µm; HS and 1.58 ± 0.35 µm; GL after 2,425 h of annealing or rather 4.25 ± 0.63 µm; HS and 2.74 ± 1.00 µm; GL after 9,808 h of annealing). Furthermore the width of gauge length sections vary stronger after longer annealing (4.25 ± 0.63 µm; HS and 2.74 ± 1.00 µm; GL after 9,808 h). Smaller zone width in case of gauge length sections in comparison to the head sections may be
caused by strain localisation and thus increased dislocation densities. This consequently potentially boosts heterogeneous precipitation due to lower nucleation energy in these areas if an oversaturated matrix is provided (1984). Detailed analyses of dislocation behaviour in the particle free zones will be discussed hereafter.
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As mentioned before the Crofer® 22 H material is strengthened by solid-solution and intermetallic Laves phase particles. The predecessor of this material is the commercially available material Crofer® 22 APU. This steel type is strengthened by solid solution only. Thus, the increased strength of the Crofer® 22 H material is mainly the result of additional particle hardening (also see (Froitzheim et al., 2008; Kuhn et al., 2011)). Therefore the observed particle-free zones are of specific interest, because in these areas the 2.5W0.57Nb0Ti material may rather behave like a solely solid-solution strengthened material. In-depth analyses of the PFZs were performed by HR-FESEM using EBSD technique for crystal (mis-)orientation mappings and HR-TEM for investigations of dislocation behaviour and dislocation densities. Band contrast images and corresponding misorientation mappings of alloy 2.5W0.57Nb0Ti in the as received condition (0 h of annealing, cf. Figure 6a+b), after 2,425 h (cf. Figure 6c-f) and after 9,808 h (cf. Figure 6g-j) of aging from the head and gauge length sections are given in Figure 6.
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Figure 6: Alloy 2.5W0.57Nb0Ti at 650 °C: a) EBSD band contrast image with grain boundaries of the as received condition (black: high-angle, red: low angle), b) misorientation mapping of the as received condition, c) EBSD band contrast image with grain boundaries image after 2,425 h of aging (head section, black: high-angle, red: low angle), d) misorientation mapping after 2,425 h (head section), e) EBSD band contrast image with grain boundaries after 2,425 h (gauge length, black: high-angle, red: low angle), f) misorientation mapping after 2,425 h (gauge length), g) grain EBSD band contrast image with grain boundaries after 9,808 h (head section, black: high-angle, red: low angle), d) misorientation mapping after 9,808 h (head section), e) EBSD band contrast image with grain boundaries after 9,808 h (gauge length, black: high-angle, red: low angle) and f) misorientation mapping after 9,808 h (gauge length) (Lopez Barrilao, 2017).
The band contrast images of the quasi stress-free head section and the stressed gauge length section show pronounced formation of low-angle grain boundaries within the PFZs. These newly formed low-
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angle grain boundaries interestingly are not occupied by particles. The reason for the formation of these boundaries is lattice deformation caused by external and/or residual stress or rather accumulation of dislocations. These low-angle boundaries are only verifiable by EBSD measurements. In general the formed boundaries in the gauge length samples show predominant orientation perpendicular to the high-angle boundaries, which are indicated in black in the left image column of Figure 6. By means of local misorientation mappings (cf. right column of Figure 6) even small differences in crystallographic lattice orientation can be displayed. These small changes in orientation are visible as misorientations of around 1° indicated as green lines. Most of the observed areas (the bcc matrix) certainly do not show strong orientation differences. Therefore most of the area is indicated in blue or rather as nearly 0° in misorientation. In comparison to the PFZ areas, greater sections of misorientation were found in the areas occupied by particles. On the contrary the misorientations in the PFZs are only marginal. Misorientations of around 1° are already visible in the as received state (cf. Figure 6a+b) without additional influence of particles, which are caused by the preceding hot-rolling process. A direct comparison of the mappings in the right column of Figure 6 indicates more pronounced crystal lattice relaxation within the PFZs. This relaxation process is only reduced due to particles. Therefore stronger misorientations arise by particles and thus affect deformation of the lattice, which might be enhanced by the presence of dislocations. These dislocations pin at particles and trigger stronger misorientations in these areas. Additional investigations of the PFZs were performed by TEM. The results provide further detail information on dislocation behaviour in these zones. Figure 7 shows an overview of TEM lamellae, which were cut from the stress-free head section and the stressed gauge length after 2,425 h of creep testing at 650 °C / 100 MPa and after 9,808 h at 70 MPa. The white lines in the left column indicate the borders of the particle-free zones. Magnifications of the areas marked by red squares are given in the right column of Figure 7. The areas indicated by dashed lines in the PFZs in the right column represent the areas where dislocation densities were calculated. The corresponding dislocation densities are given in Table 2. Density calculations only consider freely moveable and solitary dislocations. Dislocations which are already arranged in dislocation structures were neglected.
2,425 h (GL) ≈34.72
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Table 2: Dislocation densities (in 1012 m-2) calculated from the dashed marked areas in Figure 7 of the stress-free head sections and the stressed gauge lengths after 2,425 h at 100 MPa and 9,808 h at 70 MPa.
9,808 h (HS) ≈10.66
9,808 h (GL) ≈53.00
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Comparison of dislocation density calculations of the head section samples after 2,425 h and 9,808 h shows similar values, which indicates a stable state after relatively short time in the absence of stress. The influence of stress results in an increase of dislocation density. An increase in direct comparison between head section and gauge length samples was determined. After 9,808 h of creep testing at 70 MPa the density increases by approx. five times and after 2,425 h at 100 MPa an increase by a factor of three was identified. Thus, the dislocation densities of the gauge length sections apparently show an inverse dependency on the stress level, but a direct proportionality to the particle-free zones width. The dependencies of dislocation densities are given in Figure 8.
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Figure 8: Dependencies of dislocation density on particle-free zone width and stress level.
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In case of the stressed gauge length sections the examination of dislocations by TEM clearly shows a preferable orientation of the dislocations parallel to the grain boundaries (cf. Figure 7c+d and g+h). This orientation may support an accumulation of dislocations at the grain boundaries and consequently benefit void formation (Gottstein, 2007) and finally premature failure. Furthermore the preferable orientation decreases diffusion towards grain boundaries, but promotes diffusion alongside the dislocations and thus may decelerate the growth of grain boundary particles. EDX measurements of the matrix within the PFZ exhibits a tungsten content in the gauge length of approx. 1.9 wt.% in case of the 2,425 h sample tested at 100 MPa and approx. 0.65 wt.% in case of the 70 MPa sample after 9,808 h. Consequently only a third of the tungsten content remains after a four times longer testing period in the PFZs of the 70 MPa sample with enhanced dislocation density (cf. Table 2), while parts of the tungsten content conduce constantly to e.g. particle coarsening (cf. Figure 2). The tungsten contents in the PFZs in the quasi stress-free head sections do not vary significantly from the ones found in the gauge length. The head section of the 2,425 h sample tested at 100 MPa contains tungsten contents between 1.9 wt.% and 2.1 wt.%, while in case of the 9,808 h sample tested at 70 MPa tungsten contents between 0.8 wt.% and 1.0 wt.% were found. Therefor the reduction of the tungsten content is mainly a consequence of longer testing duration. The reduction in tungsten content and thus reduced solid-solution strengthening of the PFZs however might give an explanation for localized, increased deformation in the PFZs encountered in the 70 MPa sample tested for 9,808 h. In addition this influences the nucleation of Laves phase particles, which is strongly increased by enhanced dislocation densities (Hsiao et al.; Kuhn et al., 2016). Therefore the reason for smaller PFZ widths in case of gauge length sections in comparison to head sections is proposed to be a combination of preferable dislocation orientation and thus decelerated diffusion towards grain boundaries and the accompanied influence on particle nucleation and growth (cf Table 3). Final conclusions on the relations between particle-free zones, their width, dislocation densities and damage formation or rather premature failure cannot be drawn yet. More experiments are needed in order to obtain enough statistical coverage. Ongoing research focuses on the provision of satisfactory data as well as chemical composition and heat treatment optimisations to minimise the formation of PFZs.
Table 3: Mean width development and standard deviation of particle-free zones after different annealing times at 650 °C (in µm).
2h 0.45±0.11
100 h 0.63±0.18
1,000 h 0.87±0.25
2,425 h (HS) 1.72±0.38
2,425 h (GL) 1.58±0.35
9,808 h (HS) 4.25±0.63
9,808 h (GL) 2.74±0.1.00
Summary and Conclusion
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Preceding investigation of high chromium, fully ferritic steels, strengthened by intermetallic Laves phase particles contextualise particle evolution and creep behaviour (Lopez Barrilao, 2017; Lopez Barrilao et al., 2017a; Lopez Barrilao et al., 2017b; Lopez Barrilao et al., 2015). In comparison to all analysed trial alloys, the creep strongest alloy 2.5W0.57Nb0Ti (Kuhn et al., 2014) in general exhibits the tendency to maintain smaller particles during long-term high temperature application. This result was attributed to improved particle stability and/or dynamic precipitation processes causing superior creep performance. But these particle investigations did not give information about the reasons for material damage and failure. Therefore alloy 2.5W0.57Nb0Ti was analysed concerning sub-grain structure and dislocation evolution. Initial sub-grain boundaries become decorated by Laves phase particles during short-term annealing, but continually decompose until disappearance by recovery during mid-term annealing. These initial sub-grain structures are a result of the preceding hot-rolling process, with precipitates gradually occupying the cell boundaries during aging and are not to be confused with low-angle boundaries developing in particle free zones along high-angle grain boundaries during creep deformation. Comparisons of samples taken from quasi stress-free head sections and stressed gauge length sections after creep testing do not show significant differences regarding the evolution of the initial sub-grain structures. Therefore additional influence of stress on these can be assessed to be as a minor one.
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Additionally the formation of particle-free zones along high-angle grain boundaries was observed. These zones are of specific interest, because within these the material acts similar to a solely solidsolution strengthened material, what in comparison to the particle populated grain interiors is accompanied by a drop in creep strength in the long-term. The analyses show that the particle-free zones widths increase with time, but decrease with stress. Furthermore transmission electron microscopy investigations were performed to obtain detailed information about the PFZs. Increased dislocation densities as well as preferred orientation of the dislocations parallel to high-angle grain boundaries were verified. This accumulation of dislocations may benefit creep void formation and thus premature material failure. Moreover their preferred orientation may decelerate diffusion towards the grain boundaries and thus influence boundary particle growth. Additional chemical analyses show strong variations of the tungsten content depending on dislocation densities. Therefore changes in solid-solution strengthening in the PFZs can be assumed. Due to these changes in solid-solution strengthening, changes in dislocation densities and diffusion rates the particle-free zones are suspected to be a main factor contributing to creep damage and failure.
Acknowledgments The authors would like to thank B. Werner, H. Reiners, D. Liebert, A. Moser, M. Braun and W. Lange for mechanical (long-term) testing and sample annealing. Moreover the authors gratefully acknowledge the support of D. Eßer, V. Gutzeit and J. Bartsch for sample and metallographic
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preparation as well as J. Ning and W. Chen for the provided images. The support of VDM Metals GmbH with experimental material is greatly acknowledged.
References
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1984. Werkstoffkunde Stahl, Band 1: Grundlagen. Verein Deutscher Eisenhüttenleute, SpringerVerlag Berlin Heidelberg Danielsen, H.K., Hald, J., Grumsen, F.B., Somers, M.A.J., 2006. On the crystal structure of Z-phase Cr (V, Nb) N. Metallurgical and Materials Transactions A 37, 2633-2640. Froitzheim, J., Meier, G.H., Niewolak, L., Ennis, P.J., Hattendorf, H., Singheiser, L., Quadakkers, W.J., 2008. Development of high strength ferritic steel for interconnect application in SOFCs. Journal of Power Sources 178, 163-173. Gottstein, G., 2007. Physikalische Grundlagen der Materialkunde. Springer Verlag. Hald, J., 1996. Metallurgy and creep properties of new 9-12%Cr steels. Steel Research 67, 369-374. Hald, J., 2008. Microstructure and long-term creep properties of 9–12% Cr steels. International Journal of Pressure Vessels and Piping 85, 30-37. Hosoi, Y., Wade, N., Kunimitsu, S., Urita, T., 1986. Precipitation behaviour of Laves phase and it's effect on toughness of 10Cr–2Mo ferritic-martensitic steel. Journal of Nuclear Materials 141-143, 461-467. Hsiao, Z.W., Kuhn, B., Yang, S.M., Yang, L.C., Huang, S.Y., Singheiser, L., Kuo, J.C., Lin, D.Y.L.D.Y., The influence of deformation on the precipitation behavior of a ferritic stainless steel, in: LecomteBeckers, J., Dedry, O., Oakey, J., Kuhn, B. (Eds.), 10th Liege Conference : Materials for Advanced Power Engineering 2014. Kuhn, B., Asensio Jimenez, C., Niewolak, L., Hüttel, T., Beck, T., Hattendorf, H., Singheiser, L., Quadakkers, W.J., 2011. Effect of Laves phase strengthening on the mechanical properties of high Cr ferritic steels for solid oxide fuel cell interconnect application. Materials Science and Engineering: A 528, 5888-5899. Kuhn, B., Talík, M., Lopez Barrilao, J., Singheiser, L., Yamamoto, Y., Schmitz-Niederau, M., 2016. Development Status of High performance Ferritic (HiperFer) steels, 8th International Conference on Advances in Materials Technology for Fossil Power Plants, pp. 1019-1028. Kuhn, B., Talík, M., Niewolak, L., Zurek, J., Hattendorf, H., Ennis, P.J., Quadakkers, W.J., Beck, T., Singheiser, L., 2014. Development of high chromium ferritic steels strengthened by intermetallic phases. Materials Science and Engineering A, 594, 372-380. Lopez Barrilao, J., 2017. Microstructure Evolution of Laves Phase Strengthened Ferritic Steels for High Temperature Applications, Schriften des Forschungszentrums Jülich Reihe Energie & Umwelt / Energy & Environment. RWTH Aachen University, Jülich, pp. XVI, 134 pp. Lopez Barrilao, J., Kuhn, B., Wessel, E., 2016. Microstructure and Intermetallic Particle Evolution in Fully Ferritic Steels, 8th International Conference on Advances in Materials Technology for Fossil Power Plants, pp. 1029-1037. Lopez Barrilao, J., Kuhn, B., Wessel, E., 2017a. Identification, size classification and evolution of Laves phase precipitates in high chromium, fully ferritic steels. Micron 101, 221-231. Lopez Barrilao, J., Kuhn, B., Wessel, E., Talík, M., 2017b. Microstructure of intermetallic particle strengthened high-chromium fully ferritic steels. Materials science and technology 33, 1056-1064. Lopez Barrilao, J., Talík, M., Kuhn, B., Wessel, E., 2015. Microstructure Evolution of High Temperature Fully Ferritic Steels, 1st International Conference on Advanced High-Temperature Materials Technology for Sustainable and Reliable Power Engineering (123 HiMAT-2015), Sapporo, Japan. 123rd Committee of Japan Socity for the Promotion of Science, IAI, pp. 15-18. Maruyama, K., Sawada, K., Koike, J., 2001. Strengthening mechanisms of creep resistant tempered martensitic steel. ISIJ international 41, 641-653. Ning, J., 2013. Influence of chemical composition on precipitation reactions in high chromium ferritic steels. RWTH Aachen – Institute of Ferrous Metallurgy. Quadakkers, W.J., Zurek, J., 2010. Oxidation in steam and steam/hydrogen environments. Shreir's Corrosion 1, 407-456.
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Talik, M., 2016. Influence of Initial Thermomechanical Treatment on High Temperature Properties of Laves Phase Strengthened Ferritic Steels, Schriften des Forschungszentrums Jülich Reihe Energie & Umwelt / Energy & Environment. RWTH Aachen University, Jülich, pp. xxiii, 130. Viswanathan, R., Bakker, W., 2001. Materials for ultrasupercritical coal power plants—Boiler materials: Part 1. Journal of Materials Engineering and Performance 10, 81-95. Wright, I.G., Dooley, R.B., 2010. A review of the oxidation behaviour of structural alloys in steam. International Materials Reviews 55, 129-167.