- Email: [email protected]
Microstructure–mechanical property correlation in oxide dispersion strengthened 18Cr ferritic steel
Author’s Accepted Manuscript Microstructure–Mechanical Property Correlation in Oxide Dispersion Strengthened 18Cr Ferritic Steel M. Nagini, R. Vijay, Koteswararao V. Rajulapati, A.V. Reddy, G. Sundararajan www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(17)31340-0 https://doi.org/10.1016/j.msea.2017.10.023 MSA35626
To appear in: Materials Science & Engineering A Received date: 1 June 2017 Revised date: 30 September 2017 Accepted date: 4 October 2017 Cite this article as: M. Nagini, R. Vijay, Koteswararao V. Rajulapati, A.V. Reddy and G. Sundararajan, Microstructure–Mechanical Property Correlation in Oxide Dispersion Strengthened 18Cr Ferritic Steel, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2017.10.023 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.
Microstructure–Mechanical Property Correlation in Oxide Dispersion Strengthened 18Cr Ferritic Steel M. Naginia, b, R. Vijaya*, Koteswararao V. Rajulapatib, A.V. Reddya and G. Sundararajana,c a
International Advanced Research Centre for Powder metallurgy and New Materials (ARCI), Balapur, Hyderabad – 500005 (INDIA) b
School of Engineering Sciences and Technology (SEST), University of Hyderabad, Gachibowli, Hyderabad – 500046 (INDIA) c
Indian Institute of Technology, Madras, Chennai–600036 (INDIA)
ABSTRACT The tensile deformation of Oxide Dispersion Strengthened 18Cr ferritic steels (henceforth designated as ODS–18Cr steel) was studied over a temperature range of 298–1073 K. At each temperature, the influence of microstructure (grain size and dispersoid size) which could be refined progressively by increasing the milling time over the range 1 to 6 h on strength was also investigated. Oxide free 18Cr steel (NODS) provided the baseline data as compared to ODS–18Cr steel. At all the test temperatures, the flow stress of ODS–18Cr steels increased with increasing milling time or equivalently with refinement of the grains and dispersoids. The decrease in flow stress with increasing tensile test temperature was marginal up to 673 K. Beyond 673 K, the flow stress decreased rapidly. Enhanced strength of ODS steels when compared to NODS steel is due to the formation of ultra–fine grained structure along with fine dispersion of complex Y–Ti–O oxide particles. The concomitant roles of the grain size related strengthening and dispersion strengthening due to oxide particles in the strengthening of ODS–18Cr steels at all test temperatures were rationalized using root mean square superposition model.
*
Corresponding author Tel: +91 40 24443170, [email protected] (R.Vijay) 1
Key Words: Mechanical alloying, Oxide dispersion strengthened steel, Tensile properties.
1.
Introduction
Iron chromium alloys exist in ferritic condition at all temperatures when chromium content exceeds about 13 wt. % (Fig. 1). The Fe–Cr phase diagram also suggests the formation of embrittling phase containing Fe–47.2 wt. % Cr [1]. Oxide dispersion strengthened (ODS) ferritic steels are of recent origin with superior properties, such as, high temperature strength [2–7] and resistance to creep [5–10], corrosion [7, 11, 12], oxidation [7, 13, 14] and neutron irradiation [2, 7]. Therefore, they are potential candidates for a variety of high temperature applications, such as, blanket materials for fusion reactors, fuel cladding materials for Gen– IV fission reactors and blades for gas and ultra–super critical steam turbines [3, 7, 10, 15]. The improved resistance of ODS steels to high temperature tensile and creep deformation is due to the presence of thermally stable nano–sized oxide particles (Y–Ti–O) in the ductile matrix. The dispersoids [2, 7] offer resistance to dislocation motion and grain boundary sliding, and also inhibit grain growth when exposed to high temperatures. ODS ferritic steels containing high chromium (12–18 %) are being developed for improved corrosion and oxidation resistance as substitutes for relatively inferior ODS–9Cr steels [12].
ODS steels are generally processed by powder metallurgy route, consisting of mechanical alloying followed by consolidation either by hot extrusion or hot isostatic pressing (HIP). The process parameters employed in the manufacture of ODS steels, especially milling time, has a strong influence on the development of microstructure of the matrix and the size and number density of dispersoids as recently demonstrated by Nagini et al [16]. The literature available on the high temperature deformation behavior of ODS–18Cr ferritic steels is limited [17–25] when compared to the ODS–14Cr ferritic steels [25–32]. Carlan et al. and Jae Hoon Lee [17,
2
18] have evaluated the tensile properties of ODS–18Cr over a temperature range 298 to 973 K. Hadraba et al. [19] have investigated the impact properties of ODS–9Cr, ODS–14 Cr and ODS–18Cr steels as a function of processing route. Recently, Deepak Kumar et al. [20] have demonstrated that ODS–18 Cr steels can be developed through the powder forging route instead of hot extrusion. Rouffie et al. [21] have studied the effect of the thermal ageing on the tensile and impact properties of ODS–18Cr ferritic steel. Given the very limited work carried out till date on the mechanical behavior of ODS–18Cr steel, there is a clear need to understand their deformation behavior at different temperatures and also to model their mechanical behavior.
Fig.1. Fe–Cr phase diagram [1].
3
2.
Experimental Details
Inert gas atomized pre–alloyed 18Cr ferritic steel (Fe–18Cr–2.3W–0.3Ti) powder with an average particle size of 100 µm (produced at International Advanced Research Centre for Powder Metallurgy and New Materials, Hyderabad, India) was milled with and without nano yttria powder (M/s. Inframat Advanced Materials, USA, size: 30–50 nm, purity: 99.95 %) in a high energy horizontal attritor mill (Simoloyer CM–20) up to 6 h. Milling was carried out at room temperature in a water cooled stainless steel container using hardened steel balls of 5 mm under argon atmosphere. Ball to powder ratio of 7.5:1 and milling speed of 550 rpm were maintained for all milling experiments. The milled powders were filled in mild steel cans, degassed and vacuum sealed. The sealed cans were upset forged at 1323 K and hot extruded at 1423 K with an extrusion ratio of 19. The extruded rods were annealed at 1173 K for 1 h and water quenched. The finer details of the process were given in the earlier publication of the authors [16].
The chemical composition prealloyed powder as well as bulk samples was evaluated using ICPAES (JOBIN YUON FRANCE, Model: Ultima2CHR), LECO oxygen/nitrogen (Model: TC436) and LECO carbon/sulpher (Model: CS444) analyzers. The chemical compositions of pre–alloyed 18Cr steel powder, extruded and annealed oxide free 18Cr steel (NODS) and ODS18Cr steel are given in Table 1. Grain sizes of all the annealed rods were estimated from the grain boundary maps obtained by Electron Back–Scattered Diffraction (EBSD) technique available with FESEM (Hitachi, model: S4300SE/N). Grains with misorientation angle >15 are only considered for the calculation of the average size. Transmission electron microscopic (TEM) observations of annealed samples before and after tensile testing were performed using FEI Tecnai G2 200 kV (LaB6) microscope equipped with Olympus (SIS mega view III) wide angle camera. Dispersoid size, number density and 4
volume fraction calculations were carried out as per the procedures reported by the authors in their earlier publication [16].
The tensile properties of the annealed samples were evaluated as per ASTM E21 standard from room temperature (RT (298 K)) to 1073 K at a strain rate of 7.5×10–4 s–1 using Universal Testing Machine (INSTRON, Model: 4507, 200 KN). Round tensile test specimens of 3.99 mm and 22 mm gauge length were used. A constant heating rate of 10 C/min, hold time of 2 h at test temperature and furnace cooling was maintained for tensile testing of all the samples at temperatures above 298 K. For each test condition, the average results obtained on two specimens are reported. Sample/ Composition Pre–alloyed 18Cr steel powder
Fe
Cr
W
Ti
C
Total O
N
Y2O3
Excess O
bal.
17.8
2.3
0.34
<0.03
0.006
0.003
–
–
NODS
bal.
17.6
2.3
0.34
<0.03
0.05
0.012
–
–
ODS–F
bal.
17.4
2.2
0.31
<0.03
0.14
0.012
0.36
0.06
Table 1 Chemical composition (wt. %) of pre–alloyed powder, NODS and ODS–F samples.
3.
Results
3.1.
Chemical Composition
The chemical analysis indicated that the oxygen and nitrogen contents in both ODS and NODS steels increased slightly when compared to the pre–alloyed steel powder due to minor leakages during milling in argon atmosphere. Since the Y/ Ti atomic ratio continues to be 1:1 in spite of the O pick up, its importance to the overall structure and properties is negligible [33].
5
3.2.
Microstructural Features
The evolution of microstructure in ODS–18Cr steels during various processing steps (milling, upset forging, extrusion and annealing) has been studied exhaustively by the authors and published elsewhere [16]. However, for the sake of completeness, the essential microstructural features in extruded and annealed condition as a function of milling time are presented. Typical variation of TEM microstructures of ODS–18Cr steel milled for 1, 3 and 6 h of milling and of NODS obtained on the longitudinal sections is shown in Figs. 2 (a–d) respectively. The grains are elongated at 1 h milling but are more equiaxed in the case of 3 and 6 h milled ODS samples. NODS (Fig. 2 (a)) exhibits very coarse and uniform grain size. The average grain sizes (longitudinal direction) of extruded and annealed NODS and ODS steels estimated by EBSD are given in Table 2. Additionally, TEM bright field images of ODS steels milled for 1, 3 and 6 h which reveal the dispersoids are presented in Figs. 3 (a–c) respectively. The average size, number density and volume fraction of dispersoids estimated by detailed TEM examination are also included in Table 2.
6
Fig. 2. TEM images showing the grain structure of extruded+annealed samples in longitudinal direction: (a) ODS–C, (b) ODS–M, (c) ODS–F and (d) NODS.
Name
Average grain size (dg), [µm] (longitudinal)
Average dispersoid size (dp)[nm]
Number density [m–3]
Volume fraction
Inter– dispersoid spacing (S), [nm]
NODS
13 (± 6.28)
–
–
–
–
ODS–C
1.17 (± 0.27)
15 (± 11.76)
6.21x1021
0.0092
101
ODS–M
0.73 (± 0.26)
10 (± 7.03)
1.82x1022
0.0094
66
5 (± 2.35)
1.24x1023
0.0096
33
0.48 (± 0.21) 0.30 (± 0.21)* * In transverse direction ODS–F
Table 2 Average grain size, average size, number density, volume fraction and inter– dispersoid spacing of dispersoids in extruded+annealed ODS–18Cr steels.
7
Fig. 3. TEM bright field images of extruded+annealed ODS–18Cr steels: (a) ODS–C, (b) ODS–M and (c) ODS–F.
In Fig. 4, the relationship between milling times and grain size and dispersoid size using the data in Table 2 is presented in a graphical form. This figure clearly shows that increasing milling time (1 to 6 h) decreases both grain size and dispersoid size simultaneously. Henceforth, ODS–18 Cr steel milled for 1, 3 and 6 h will be designated as ODS–C, ODS–M and ODS–F with the letters C, M and F denoting coarse, medium and fine microstructures.
8
Fig. 4. The contribution of grain size and dispersed size obtained for 1, 3 and 6 h milling.
To assess the combined effect of strain and temperature on the extent of coarsening of the grains and dispersoids during the tensile test, microstructural examination was carried out on ODS–F sample after tensile deformation at 298 and 1073 K. The requisite sample in transverse direction was extracted from around 1 mm below the fracture surface for the above purpose. The relevant microstructural features showing the grain structures and dispersoids of ODS–F in annealed and deformed (298 and 1073 K) conditions are presented in Figs. 5 (a– c)–(d–f) respectively. The corresponding quantified data of grain and dispersoid sizes of ODS–F sample are given in Table 3. The data presented clearly indicates that neither the grains nor the dispersoids exhibit any coarsening even when exposed to a temperature of 1073 K during the tensile test. It is well established that the dispersoids are formed by the precipitation of Y–Ti–O complexes during hot consolidation either from the ferrite containing Y, Ti and O in solid solution or from nano clusters in as milled condition [34]. All 9
the solutes are expended in the formation of stable Y–Ti–O complex oxides during high temperature exposure, leaving little or no Y left in solid solution to promote growth of dispersoids. Further, because of the stability of Y–Ti–O complex oxide dispersoids, coarsening does not occur up to 1200 K even if assisted by strain. This observation implies that the grain size, dispersoid size and their number density can be assumed to be invariant with test temperature while modeling the yield strength behavior of ODS–C, ODS–M and ODS–F steels.
Fig. 5. TEM images showing the grain structures and dispersoids of ODS–F in (a and d annealed condition, (b and e) deformed at RT (298 K) and (c and f) deformed at 1073 K, respectively.
Sample name
ODS–F
Sample condition
Average Average grain size (dg), [µm] dispersoid Size (dp) (transverse direction) [nm]
Annealed
0.32 (± 0.18)
4.9 (± 2.35)
Tensile tested at 298 K
0.3 (± 0.14)
4.5 (± 1.79)
Tensile tested at 1073 K
0.31 (± 0.17)
4.6 (± 1.93)
Table 3 Average grain size (transverse) and dispersoid size in ODS–F in annealed sample and in deformed condition at RT (298 K) and 1073 K.
10
3.3.
Tensile Properties
Tensile tests were carried out on ODS–18Cr and NODS steels from 298 to 1073 K to assess the effect of microstructure, resulting from varying milling times, on strength and ductility. The engineering stress–strain curves of NODS and ODS–18Cr steels at different test temperatures are shown in Figs. 6 (a–d) respectively. The variation of yield strength (y) and ultimate tensile strengths (u) of ODS–C, ODS–M and ODS–F along with NODS with temperature is shown in Figs. 7 (a–b) respectively. The uniform (eu) and total fracture strains (et) for NODS and ODS–18Cr steels are shown in Figs. 8 (a–b) respectively. From the tensile data presented in Figs. 6 to 8, the following conclusions can be drawn. (a) Fig. 6 indicates that at low test temperatures (673 K and below), the shape of the tensile curves of both ODS–18Cr and NODS steels changed from characteristic yielding to work hardening to softening (necking) and finally fracture. At temperatures above 673 K, the work hardening phase was absent. (b) The engineering stress–strain curves (Fig. 6) of all the alloys (including NODS) indicate that the decrease in yield strength and ultimate tensile strength (UTS) with increasing temperatures is marginal up to 673 K but more dramatic at higher temperatures. Such behavior is clearly seen in Figs. 7 as well. (c) The effect of progressive refining of the microstructure (i.e, ODS–C to ODS–M to ODS–F) is to increase the yield strength and UTS at all temperatures–even at 800 C (1073 K). Such behavior can be attributed to the thermal stability of the grains and dispersoids even at 1073 K as observed in Fig. 5. (d) The ductility values of ODS alloys (Fig. 8 (a)), covering the temperature range 298 to 1073 K, lie in the range 20 to 35 % which is quite acceptable. NODS alloy exhibits higher ductility as compared to ODS alloys. The difference in fracture strain between
11
NODS and ODS alloys is due to the drastic reduction in strain hardening in ODS alloys which in turn is caused by the fine grained structure. The unusually large fracture strain in NODS at above 873 K is attributable to near zero strain hardening exponent and favourable temperature, grain size and strain rate hardening. The ODS– 18Cr steels on the other hand showed decrease in fracture strain due to the stability of microstructure which promotes inhomogeneous deformation and pinning of grain boundaries by oxide particles inhibiting grain boundary sliding even at high temperatures. (e) In all alloys (including NODS), the uniform strain decreases progressively with increasing test temperature (see Fig. 8 (b)). NODS exhibit the highest uniform strain values and ODS–F the lowest.
Fig. 6. Engineering stress–strain curves of 18Cr and ODS–18Cr steels at different test temperatures: (a) NODS, (b) ODS–C, (c) ODS–M and (d) ODS–F.
12
Fig. 7. Variation of (a) y and (b) u of 18Cr and ODS–18Cr steels with test temperature.
Fig. 8. Variation of ductility indices of (a) eu and (b) et of 18Cr and ODS–18Cr steels with test temperature.
4.
Discussion
4.1.
Microstructural Changes with Milling Time
An unique feature of the present work is the demonstration that using milling time as one of the process variables, one can obtain a range of microstructures (coarse to medium to fine) in the same ODS–18Cr alloy. The events leading to the refinement of microstructure with increasing milling time are reasonably well understood. Many studies [16, 35–37] have suggested that, Y2O3 particles added to the pre–alloyed 18Cr steel powder fragment progressively with increasing milling time and ultimately dissolve in the matrix. Simultaneously, nano clusters of Y–Ti–O start forming during the final stages of milling. These Y, Ti and O nano clusters transform to Y–Ti–O nano dispersoids during subsequent 13
hot extrusion and annealing. Higher the milling time, finer and more numerous are the nanoclusters and hence finer are the dispersoids even after extrusion and annealing. Thus, the effect of milling time persists on the microstructure even after subsequent extrusion and annealing. The sequence of events with regard to grain size is similar. With increasing milling time, the grain size reduces in size but some grain growth occurs during subsequent hot extrusion and annealing. However, the influence of milling time still persists. It is also clear that the nano dispersoids are very effective in pinning the grain boundaries resulting in extremely fine grain size in ODS–18Cr steel (0.48 to 1.17 m) as compared to NODS wherein the final grain size is an order of magnitude higher at 13 m.
4.2.
Influence of Microstructure and Temperature on Yield Strength
The strong influence of microstructure on yield strength at all temperatures (298 to 1073 K) is evident from Table 4, wherein the yield strength of ODS–C, ODS–M and ODS–F is presented at various temperatures. The increase in yield strength due to refinement of microstructure (σy(F)–σy(C), Table 4) is 351 MPa at 298 K and decreases with increasing temperature to 116 MPa at 1073 K. Looked another way, the ratio of yield strength of ODS–F to ODS–C (see Table 4) is around 1.76 at room temperature (298 K) and increases to 2.0 at 1073 K. In contrast, the ratio of yield strength at 1073 to 298 K, in the case of ODS–C, ODS– M and ODS–F alloys, provided in the bottom row of Table 4, is reasonably constant (0.25 to 0.28). Thus, the influence of microstructure persists even at the highest test temperature of 1073 K. All these aspects of microstructure–yield strength correlation will be understood on the basis of a model for yield strength proposed in the next section.
14
σy [MPa] ∆σy [MPa] σy (ODS–F) Temperature [K] C M F σy(F)–σy(C) σy (ODS–C) 298 459 701 810 351 1.76 673 377 549 662 285 1.76 873 244 375 453 209 1.86 1073 116 192 232 116 2.00 σy (1073 K) 0.25 0.27 0.28 – – σy (298 K) Table 4 The yield strength of ODS–18Cr steels as a function of temperature and microstructure (C, M and F).
4.3.
Modeling of the Yield Strength at Room and Elevated Temperatures
The possible strengthening mechanisms contributing to the yield strength of ODS–18Cr steels include Peierls–Nabarro lattice resistance stress, solid solution strengthening, dislocation forest hardening, Hall–Petch grain boundary strengthening and dispersion strengthening due to Orowan mechanism.
In the present study, the yield strength of not only the ODS–18Cr steels (C, M and F) but also of NODS steel, which does not have the dispersoids, have been experimentally obtained. Thus, the difference in yield stress between ODS–18Cr steel and NODS steel is largely due to Hall–Petch grain boundary strengthening and dispersion strengthening due to Orowan mechanism. The extent of strengthening due to Peierls–Nabarro lattice resistance, solid solution strengthening and forest dislocation hardening should be the same in both NODS and ODS–18Cr steel. This is especially true because the solid solution strengthening elements like W, Cr do not get incorporated in the dispersoid and thus are available for solid solution strengthening to the same extent in NODS and ODS–18Cr steels. Additionally, since both NODS and ODS–18Cr steel have been annealed at 1173 K after extrusion, it is likely that the dislocation density exhibits low and similar values in both materials. However, NODS also includes some contribution from Hall–Petch strengthening since it has an average grain size 15
of 13 m. Thus, the matrix strength (m), which includes only the contributions from Peierls lattice stress, solid solution strengthening and forest dislocation hardening is now given as,
m y ,NODS H P,NODS
(1)
Finally, the yield strength of ODS–18Cr steels is obtained as,
y ,ODS m H P Or
(2)
In Eq. (2), ∆σH–P represents the yield strength increment due to grain boundary strengthening (Eq. (3)) and ∆σOr is the yield strength increment due to dispersion strengthening (Eqs. (4) and (5)). H P K H P d g1/ 2 ,
(3)
Gb d p Or A ln , S 2b
(4)
S
2 3 f 2
dp
1/ 2
22 3
1/ 2
,
(5)
To calculate the yield strength at all test temperatures (298 to 1073 K), it is important to understand the influence of temperature on the various parameters used in Eqs. (3)–(5). We have already demonstrated that grain size and dispersoid size does not vary even up to 1073 K (see Table 3 and Fig. 5). Thus, the values of dg, dp, S and f listed in Table 5 are valid at all temperatures. The shear modulus of the alloys decreases with increasing temperature as per Eq. (6) given below [38].
G(GPa) 93.2(1 4.368 104 T )
(6)
Another parameter which changes with temperature is the Hall–Petch constant (KH–P) and for the present alloy, the data of Hall–Petch constant vs. temperature provided by Kim et al. [38] in the case of ODS–14Cr steel has been utilized. Lastly, the matrix strength (σm) also varies with temperature. However, the yield strength of NODS alloy as a function of temperature
16
has been obtained in the present study and this data directly can be utilized for determining σm (as per Eq. (1)) at various temperatures.
Parameter
Values chosen NODS
ODS–C
ODS–M
ODS–F
Ref.
A
Numerical constant
–
0.3
0.3
0.3
45
b
Burgers vector [nm]
0.25
0.25
0.25
0.25
45
KH–P
Hall–Petch constant [MPa.m1/2]
338
338
338
338
38
dg
Grain size [µm]
13.0
1.17
0.73
0.48
Present study
dp
Dispersoid size [nm]
–
15
10
5
Present study
S
Inter–dispersoid spacing [nm]
–
101
66
33
Present study
f
Volume fraction of dispersoids
–
0.0092
0.0094
0.0096
Present study
Table 5 The values chosen for the various parameters appearing in the model (Eqs. (1)–(5)) for NODS and ODS–18Cr steels and the reference source for the data.
On the basis of the equations and values for the various parameters detailed above, it is now possible to calculate the variation of yield strength of ODS–Cr steels (σy,ODS; Eq. (2)) with temperature and as a function of microstructure (C, M and F). At this stage it is important to note that Eq. (2) used to calculate σy,ODS assumes that the yield stress increments due to various strengthening mechanisms can be linearly summed up and this in turn implies that there is no influence of one strengthening mechanism on the other. The prediction of the linear summation model as a function of both temperature and microstructure is presented in Fig. 9. It is clear that linear summation predicts yield strength values which are considerably higher than the experimental values at all test temperatures and at different microstructures. This is not surprising since dislocations participate in overcoming many of the obstacles present in the ODS alloys and availability of dislocations for participating in any one
17
strengthening mechanism is limited and dependent on other strengthening mechanisms which are also operating at the same time.
A number of researchers have analyzed the problem of summation of various strengthening mechanisms [39–47]. Based on extensive simulation of dislocation motion in the presence of multiple obstacles, Brown and Ham [39] have concluded that linear summation is a poor approximation except when a few strong obstacles are dispersed among a large number of weak obstacles and that quadratic summation has more validity. A generalized version of the non–linear summation is given by Eq. (7) given below.
yk 1k 2k 3k
(7)
In Eq. (7), σ1, σ2 and σ3 represent 3 different strengthening mechanisms. The value of parameter ‘k’ defines the nature of interaction between the 3 strengthening mechanisms. A value of 1 for k, transforms Eq. (7) to one representing linear summation like Eq. (2). On the other hand, k=2 represents root mean square or quadratic summation and implies that the strengthening mechanisms have synergistic interaction.
Koppenaal and Kulhman–Wilsdorf [40] have proposed the quadratic summation (k=2) model which has been supported by subsequent statistical calculations and simulations. Simulation of combined effect of Orowan and forest hardening by Queyreau et al. [41] indicated a value of 2 for k as most suitable. Lagerpusch et al. [42] investigated experimentally the superposition of solid solution strengthening and dispersion strengthening in a Cu–Au–SiO2 alloy and concluded that k=1.8 gives the best fit of the experimental data. It has been shown that when the obstacle strengths and number density are comparable, a value of 2 for k is the most appropriate [43]. Hansen [44] has demonstrated that superposition of grain boundary and particle strengthening is best explained by assuming a value of 2 for k. Nembach [45] in 18
an interesting study of the interaction between grain boundaries and precipitates in a nickel base superalloy, observed that when the precipitate size is less than 20 nm, most of the dislocations were piled up at grain boundaries and only a few Orowan loops were observed. In contrast, when the precipitate size is greater than 25 nm, only Orowan loops were seen. Thus, the interaction between two strengthening mechanisms can be quite complex.
Given the above scenario and also since the obstacles in the present ODS–18Cr alloys (grain boundary and dispersoids) have comparable strength, a k value of 2 is the most appropriate. Thus, the equation for yield strength of ODS alloy can be written down as, 2 y ,ODS ( rmss) ( m2 H2 P Or )
1
2
(8)
In the above equation, the yield strength is referred to as σy,ODS (rmss) wherein the term rmss stands for root mean square superposition.
A few investigators [46, 47] have suggested that the resistance experienced by dislocations to Peierls lattice stress, solid solution and forest dislocations can be treated as a friction stress impeding dislocation motion and in such a case linear superposition (k=1) is appropriate for some of the strengthening mechanisms. The above aspect can be taken into account in the present alloy by adding matrix strength (σm) to the root mean square of grain boundary and dispersion strengthening as indicated below. 2 y ,ODS ( prmss) m ( H2 P Or )
1
2
(9)
In Eq. (9), the yield strength is referred to as σy,ODS (prmss) wherein the term prmss stands for part root mean square summation.
19
Finally, the variation of the calculated yield strength with temperature and microstructure assuming rms summation and part rms summation are presented in Fig. 9. This figure clearly suggests that rms superposition predicts the yield strength most accurately as compared to the
predictions assuming part rms or linear superposition. Fig. 9. Comparison of the experimentally obtained yield strength as a function of temperature of ODS–18Cr steel for three microstructural conditions (ODS–C, ODS–M & ODS–F) with the predictions of yield strength assuming linear, root mean square (rms) and partial rms superposition.
Now that the ability of the model to predict the yield strength as a function of temperature and microstructure has been validated, it is possible to assess the relative contributions of grain boundary strengthening (∆σH–P) and dispersion strengthening (∆σOr). The ratio of ∆σH–P to ∆σOr is presented in Fig. 10 as a function of temperature and for three microstructures (ODS–C, ODS–M and ODS–F). This figure clearly indicates that grain boundary strengthening contribution to yield strength is greater than that of dispersion strengthening up 20
to a temperature of 873 K. At the highest temperature of 1073 K the contribution of dispersion strengthening is substantially higher than grain boundary strengthening. This can be explained on the basis of weakening of the grain boundaries at temperature higher than 50 % of the melting point. In contrast, the variation of dispersion strengthening follows the temperature dependence of the shear modulus since the dispersoids are stable even at this temperature.
Fig. 10. The variation of the ratio ∆σH–P / ∆σOr with temperature for ODS–18Cr with coarse (C), medium (M) and fine (F) microstructure.
5.
Conclusions
The tensile deformation behavior of ODS–18Cr ferritic steels along with 18Cr (NODS) steel at indifferent test temperatures have been investigated. The salient features from the present study are given below: The strength of ODS–18Cr steels increased at all the test temperatures due to the refinement of dispersoids and grains. 21
All ODS–18Cr steels exhibited superior strength values over the whole range of test temperatures compared to NODS steel and the strength of both the steels decreased with temperature. The ductility indices of ODS–18Cr steels showed marginal decrease with refinement of microstructure over the test temperature range and were lower than those of NODS steel. With increasing test temperature, NODS and ODS–18Cr steels exhibited continuous decrease in uniform elongation, but the fracture strain remained nearly same. The strength values, predicted by root mean square superposition model at room and higher temperatures compare well with experimental values.
Acknowledgements The authors thank Dr. G. Ravi Chandra, Mr. M. Ramakrishna and Mr. G.V.R. Reddy for the help for EBSD, TEM and SEM studies. They also thank Dr. K. Satya Prasad for his help in carrying out microstructural examination. The authors express their gratitude to IGCAR, Kalpakkam for funding (No. IGC/MMG/MMD/ODS/01/2010) the work and NFC, Hyderabad for carrying out hot extrusion.
References [1]
ASM Handbook Vol. 3: Alloy Phase Diagrams, ASM International, Materials Park, Ohio, 1998, 44073–0002.
[2]
G.R. Odette, M.J. Alinger, B.D Wirth, Recent developments in irradiation–resistant steels, Annu. Rev. Mater. Res. 38 (2008) 471–503.
[3]
S. Ukai, M. Fujiwara, Perspective of ODS alloys application in nuclear environments, J. Nucl. Mater. 307–311 (2002) 749–757.
[4]
A. Ramar, P. Spatig, R. Schaublin, Analysis of high temperature deformation mechanism in ODS EUROFER97 alloy, J. Nucl. Mater. 382 (2008) 210–216. 22
[5]
S. Ohtsuka, S. Ukai, M. Fujiwara, T. Kaito, T. Narita, Improvement of creep strength of 9CrODS martensitic steel by controlling excess oxygen and titanium concentrations, Mater. Trans. 46 (2005) 487–492.
[6]
R.L. Klueh, J.P. Shingledecker, R.W. Swindeman, D.T. Hoelzer, Oxide dispersion– strengthened steels: A comparison of some commercial and experimental alloys, J. Nucl. Mater. 341 (2005) 103–114.
[7]
S. Ukai, Oxide dispersion strengthened steels, Comp. Nucl. Mater. 4 (2012) 241–271.
[8]
S. Ukai, S. Mizuta, T. Yoshitake, T. Okuda, M. Fujiwara, S. Hagi, T. Kobayashi, Tube manufacturing and characterization of oxide dispersion strengthened ferritic steels, J. Nucl. Mater. 283–287 (2000) 702–706.
[9]
A. Alamo, V. Lambard, X. Averty, M.H. Mathon, Assessment of ODS–14%Cr ferritic alloy for high temperature applications, J. Nucl. Mater. 329–333 (2004) 333–337.
[10]
T. Jayakumar, M.D. Mathew, K. Laha, High temperature materials for nuclear fast fission and fusion reactors and advanced fossil power plants, Procedia Eng. 55 (2013) 259–270.
[11]
H.S. Cho, A. Kimura, S. Ukai, M. Fujiwara, Corrosion properties of oxide dispersion strengthened steels in super–critical water environment, J. Nucl. Mater. 329–333 (2004) 387–391.
[12]
A. Kimura, H.S. Cho, N. Toda, R. Kasada, K. Yutani, H. Kishimoto, N. Iwata, S. Ukai, M. Fujiwara, High burn up fuel cladding materials R&D for advanced nuclear systems: Nano–sized oxide dispersion strengthening steels, J. Nucl. Sci. Technol. 44 (2007) 323–328.
[13]
T. Kaito, T. Narita, S. Ukai, Y. Matsuda, High temperature oxidation behavior of ODS steels, J. Nucl. Mater. 329–333 (2004) 388–1392.
[14]
J.H. Kim, K.M. Kim, T.S Byun, D.W. Lee, C.H. Park, High temperature oxidation behavior of nano–structured ferritic oxide dispersion strengthened alloys, Thermochem. Acta 579 (2014) 1–8.
[15]
C. Capdevila, M. Serrano, M. Campos, High strength oxide dispersion strengthened steels: fundamentals and applications, Mater. Sci. Technol. 30 (2014) 1655–1657.
[16]
M. Nagini, R. Vijay, K.V. Ralulapati, K. Bhanu Sankara Rao, M. Ramakrishna, A.V. Reddy, G. Sundararajan, Effect of process parameters on microstructure and hardness of oxide dispersion strengthened 18Cr ferritic steel, Metall. Mater. Trans. A, 47A (2016) 4197–4209.
[17]
J.H. Lee, Development of oxide dispersion strengthened ferritic steels with and without aluminum, Front. Energy 6 (2012) 29–34.
[18]
Y. de Carlan, J.L. Bechade, P. Dubuisson, J.L Seran, P. Billot, A. Bougault, T. Cozzika, S. Doriot, D. Hamon, J. Henry, M. Ratti, N. Lochet, D. Nunes, P. Olier, T. Leblond, M.H. Mathon, CEA developments of new ferritic ODS alloys for nuclear applications, , J. Nucl. Mater. 386–388 (2009) 430–432. 23
[19]
H. Hadraba, B. Fournier, L. Stratil, J. Malaplate, A.L. Rouffie, P. Wident, L. Ziolek, J.L. Bechade, Influence of microstructure on impact properties of 9–18%Cr ODS steels for fusion/fission applications, J. Nucl. Mater. 411 (2011) 112–118.
[20]
Deepak Kumar, Ujjwal Prakash, Vikram V. Dabhade, K. Laha, T. Sakthivel, Development of oxide dispersion strengthened (ODS) ferritic steel through powder forging, J. Mater. Eng. Perform. 26 (2017) 1817–1824.
[21]
A.L. Rouffie, J. Crepin, M. Sennour, B. Tanguy, A. Pineau, D. Hamon, P. Wident, S. Vincent, V. Garat, B. Fournier, Effect of the thermal ageing on the tensile and impact properties of a 18% Cr ferritic steel, J. Nucl. Mater. 445 (2014) 37–42.
[22]
S. Li, Z. Zhou, M. Li, M. Wang, G. Zhang, Microstructure characterization and tensile properties of 18Cr–4Al oxide dispersion strengthened ferritic steel, J. Alloys Comp. 648 (2015) 39–45.
[23]
P. Dubuisson, Y. de Carlan, V. Garat, M. Blat, ODS Ferritic/martensitic alloys for sodium fast reactor fuel pin cladding, J. Nucl. Mater. 428 (2012) 6–12.
[24]
S. Li, Z. Zhou, J. Jang, M. Wang, H. Hu, H. Sun, L. Zou, G. Zhang, L. Zhang, The influence of Cr content on the mechanical properties of ODS ferritic steels, J. Nucl. Mater. 455 (2014) 194–200.
[25]
B. Fournier, A. Steckmeyer, A.L. Rouffie, J. Malaplate, J. Garnier, M. Ratti, P. Wident, L. Ziolek, I. Tournie, V. Rabeau, J.M. Gentzbittel, T. Kruml, I. Kubena, Mechanical behaviour of ferritic ODS steels–Temperature dependency and anisotropy, J. Nucl. Mater. 430 (2012) 142–149.
[26]
J.H. Kim, T.S. Byun, J.H. Lee, J.Y. Min, S.W. Kim, C.H. Park, B.H. Lee, Effects of processing condition on the microstructural and tensile properties of 14Cr–based oxide dispersion strengthened alloys, J. Nucl. Mater. 449 (2014) 300–307.
[27]
P. He, R. Lindau, A. Moeslang, H.R.Z. Sandim, The influence of thermomechanical processing on the microstructure and mechanical properties of 13.5Cr ODS steels, Fusion Eng. Des. 88 (2013) 2448–2452.
[28]
R. Didomizio, S. Huang, L. Dial, J. Ilavsky, M. Larsen, An assessment of milling time on the structure and properties of a nanostructured ferritic alloy (NFA), Metall. Mater. Trans. A, 45 A (2014) 5409–5418.
[29]
A. Garcia–Junceda, M. Hernandez–Mayoral, M. Serrano, Influence of the microstructure on the tensile and impact properties of a 14Cr ODS steel bar, Mater. Sci. Eng. A, 556 (2012) 696–703.
[30]
M. Serrano, M.H. Mayoral, A.G. Junceda, Microstructural anisotropy effect on the mechanical properties of a 14Cr ODS steel, J. Nucl. Mater. 428 (2012) 103–109.
[31]
A. Steckmeyer, M. Praud, B. Fournier, J. Malaplate, J. Garnier, J.L. Bechade, I. Tournie, A. Tancray, A. Bougault, P. Bonnaillie, Tensile properties and deformation mechanisms of a 14Cr ODS ferritic steel, J. Nucl. Mater. 405 (2010) 95–100.
24
[32]
M. Praud, F. Mompiou, J. Malaplate, D. Caillard, J. Garnier, A. Steckmeyer, B. Fournier, Study of the deformation mechanisms in a Fe–14% Cr ODS alloy, J. Nucl. Mater. 428 (2012) 90–97.
[33]
S. Ukai, S. Mizuta, M. Fujiwara, T. Okuda, T. Kobyashi, Development of 9Cr–ODS martensitic steel claddings for fuel pins by means of ferrite to austenite phase transformation, J. Nucl. Sci. Technol. 39 (2002) 778–788.
[34]
P. Miao, G.R. Odette, T. Yamamoto, M. Alinger, D. Klingensmith, Thermal stability of nano–structured ferritic alloy, J. Nucl. Mater. 377 (2008) 59–64.
[35]
M. Ratti, D. Leuvrey, M.H. Mathon, Y. de Carlan, Influence of titanium on nano– cluster (Y, Ti, O) stability in ODS ferritic materials, J. Nucl. Mater. 386–388 (2009) 540–543.
[36]
C. Cayron, A. Montani, D. Venet, Y. de Carlan, Identification of new phases in annealed Fe–18CrWTi ODS powders, J. Nucl. Mater. 399 (2010) 219–224.
[37]
Olena Kalokhtina, B. Radiguet, Y. de Carlan, P. Pareige, Study of nano–cluster formation in Fe–18Cr ODS ferritic steel by atom probe tomography, Mater. Res. Soc. Symp. Proc. 1264 (2010) 1264–BB03–04.
[38]
J.H. Kim, T.S Byun, D.T. Hoelzer, C.H. Park, J.T. Yeom, J.K. Hong, Temperature dependence of strengthening mechanisms in the nanostructured ferritic alloy 14YWT: Part II–Mechanistic models and predictions, Mater. Sci. Eng. A, 559 (2013) 111–118.
[39]
L.M. Brown, R.K. Ham, Dislocation–particle interactions, in: A. Kelly, R.B. Nicholson (Eds), Strengthening methods in solids, Applied Science Publishers, London, 1971, 9–135.
[40]
T.J. Koppenaal, D. Kuhlmann–Wilsdorf, The effect of prestressing on the strength of neutron–irradiated copper single crystals, Appl. Phys. Lett. 4 (1964) 59–61.
[41]
S. Queyreau, G. Monnet, B. Devincre, Orowan strengthening and forest hardening superposition examined by dislocation dynamics simulations, Acta Mater. 58 (2010) 5586–5595.
[42]
U. Lagerpusch, V. Mohles, E. Nembach, On the additivity of solid solution and dispersion strengthening, Mater. Sci. Eng. A 319–321 (2001) 176–178.
[43]
Y. Dong, T. Nogaret, W.A. Curtin, Scaling of dislocation strengthening by multiple obstacle types, Metall. Mater. Trans. A 41 A (2010) 1954–1960.
[44]
N. Hansen, Super position of Grain–boundary strengthening and particle strengthening in aluminium, Scr. Mater. 5 (1971) 417–420.
[45]
E. Nembach, Synergistic effects in the super position of strengthening mechanisms, Acta Mater. 40 (1992) 3325–3330.
[46]
M. Dade, J. Malaplate, J. Garnier, F. De Geuser, F. Barcelo, P.Wident, A. Deschamps, Acta Materialia, Influence of microstructural parameters on the
25
mechanical properties of oxide dispersion strengthened Fe–14Cr steels, Acta Mater. 127 (2017) 165–177. [47]
J. Lu, O. Omotoso, J.B. Wiskel, D.G. Ivery, Strengthening mechanisms and their relative contributions to the yield strength of micro alloyed steels, Metall. Mater. Trans. A 43 A (2012) 3043–3061.
26
TABLES Table 1.
Chemical composition (wt. %) of pre–alloyed powder, NODS and ODS–F samples.
Table 2.
Average grain size, average size, number density, volume fraction and inter– dispersoid spacing of dispersoids in extruded+annealed ODS–18Cr steels.
Table 3.
Average grain size (transverse) and dispersoid size in ODS–F in annealed sample and in deformed condition at RT (298 K) and 1073 K.
Table 4.
The yield strength of ODS–18Cr steels as a function of temperature and microstructure (C, M and F).
Table 5.
The values chosen for the various parameters appearing in the model (Eqs. (1)–(5)) for NODS and ODS–18Cr steels and the reference source for the data.
FIGURE CAPTIONS Fig. 1. Fe-Cr phase diagram. Fig. 2. TEM images showing the grain structure of extruded+annealed samples in longitudinal direction: (a) ODS–C, (b) ODS–M, (c) ODS–F and (d) NODS. Fig. 3. TEM bright field images of extruded+annealed ODS–18Cr steels: (a) ODS–C, (b) ODS–M, (c) ODS–F Fig. 4. The contribution of grain size and dispersoid size obtained for 1, 3 and 6 h milling. Fig. 5. TEM images showing the grain structures and dispersoids of ODS–F in (a and d) annealed condition, (b and e) deformed at RT (298 K) and (c and f) deformed at 1073 K, respectively. Fig. 6. Engineering stress–strain curves of 18Cr and ODS–18Cr steels at different test temperatures: (a) NODS, (b) ODS–C, (c) ODS–M and (d) ODS–F. Fig. 7. Variation of (a) y and (b) u of 18Cr and ODS–18Cr steels with test temperature. Fig. 8. Variation of ductility indices of (a) eu and (b) et of 18Cr and ODS–18Cr steels with test temperature. Fig. 9. Comparison of the experimentally obtained yield strength as a function of temperature of ODS–18Cr steel for three microstructural conditions (ODS–C, ODS– M and ODS–F) with the predictions of yield strength assuming linear, root mean square (rms) and partial rms superposition. Fig.10. The variation of the ratio ∆σH–P /∆σOr with temperature for ODS–18Cr steel with coarse (C), medium (M) and fine (F) microstructure. 27
Graphical abstract
High lights:
Tensile deformation behavior of nano oxide dispersoid strengthened 18Cr ferritic steels was studied.
The grain size of ODS–18Cr steels progressively decreased with the size of dispersoids.
Presence of nano sized dispersoids enhanced the strength at all temperatures when compared to dispersoid free 18Cr steel.
The flow strength of ODS–18Cr steels also increased with refinement of microstructure (grain and dispersoid size) at all test temperature.
Microstructure–mechanical property correlations were developed using root mean square superposition model.
28
PII: DOI: Reference:
S0921-5093(17)31340-0 https://doi.org/10.1016/j.msea.2017.10.023 MSA35626
To appear in: Materials Science & Engineering A Received date: 1 June 2017 Revised date: 30 September 2017 Accepted date: 4 October 2017 Cite this article as: M. Nagini, R. Vijay, Koteswararao V. Rajulapati, A.V. Reddy and G. Sundararajan, Microstructure–Mechanical Property Correlation in Oxide Dispersion Strengthened 18Cr Ferritic Steel, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2017.10.023 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.
Microstructure–Mechanical Property Correlation in Oxide Dispersion Strengthened 18Cr Ferritic Steel M. Naginia, b, R. Vijaya*, Koteswararao V. Rajulapatib, A.V. Reddya and G. Sundararajana,c a
International Advanced Research Centre for Powder metallurgy and New Materials (ARCI), Balapur, Hyderabad – 500005 (INDIA) b
School of Engineering Sciences and Technology (SEST), University of Hyderabad, Gachibowli, Hyderabad – 500046 (INDIA) c
Indian Institute of Technology, Madras, Chennai–600036 (INDIA)
ABSTRACT The tensile deformation of Oxide Dispersion Strengthened 18Cr ferritic steels (henceforth designated as ODS–18Cr steel) was studied over a temperature range of 298–1073 K. At each temperature, the influence of microstructure (grain size and dispersoid size) which could be refined progressively by increasing the milling time over the range 1 to 6 h on strength was also investigated. Oxide free 18Cr steel (NODS) provided the baseline data as compared to ODS–18Cr steel. At all the test temperatures, the flow stress of ODS–18Cr steels increased with increasing milling time or equivalently with refinement of the grains and dispersoids. The decrease in flow stress with increasing tensile test temperature was marginal up to 673 K. Beyond 673 K, the flow stress decreased rapidly. Enhanced strength of ODS steels when compared to NODS steel is due to the formation of ultra–fine grained structure along with fine dispersion of complex Y–Ti–O oxide particles. The concomitant roles of the grain size related strengthening and dispersion strengthening due to oxide particles in the strengthening of ODS–18Cr steels at all test temperatures were rationalized using root mean square superposition model.
*
Corresponding author Tel: +91 40 24443170, [email protected] (R.Vijay) 1
Key Words: Mechanical alloying, Oxide dispersion strengthened steel, Tensile properties.
1.
Introduction
Iron chromium alloys exist in ferritic condition at all temperatures when chromium content exceeds about 13 wt. % (Fig. 1). The Fe–Cr phase diagram also suggests the formation of embrittling phase containing Fe–47.2 wt. % Cr [1]. Oxide dispersion strengthened (ODS) ferritic steels are of recent origin with superior properties, such as, high temperature strength [2–7] and resistance to creep [5–10], corrosion [7, 11, 12], oxidation [7, 13, 14] and neutron irradiation [2, 7]. Therefore, they are potential candidates for a variety of high temperature applications, such as, blanket materials for fusion reactors, fuel cladding materials for Gen– IV fission reactors and blades for gas and ultra–super critical steam turbines [3, 7, 10, 15]. The improved resistance of ODS steels to high temperature tensile and creep deformation is due to the presence of thermally stable nano–sized oxide particles (Y–Ti–O) in the ductile matrix. The dispersoids [2, 7] offer resistance to dislocation motion and grain boundary sliding, and also inhibit grain growth when exposed to high temperatures. ODS ferritic steels containing high chromium (12–18 %) are being developed for improved corrosion and oxidation resistance as substitutes for relatively inferior ODS–9Cr steels [12].
ODS steels are generally processed by powder metallurgy route, consisting of mechanical alloying followed by consolidation either by hot extrusion or hot isostatic pressing (HIP). The process parameters employed in the manufacture of ODS steels, especially milling time, has a strong influence on the development of microstructure of the matrix and the size and number density of dispersoids as recently demonstrated by Nagini et al [16]. The literature available on the high temperature deformation behavior of ODS–18Cr ferritic steels is limited [17–25] when compared to the ODS–14Cr ferritic steels [25–32]. Carlan et al. and Jae Hoon Lee [17,
2
18] have evaluated the tensile properties of ODS–18Cr over a temperature range 298 to 973 K. Hadraba et al. [19] have investigated the impact properties of ODS–9Cr, ODS–14 Cr and ODS–18Cr steels as a function of processing route. Recently, Deepak Kumar et al. [20] have demonstrated that ODS–18 Cr steels can be developed through the powder forging route instead of hot extrusion. Rouffie et al. [21] have studied the effect of the thermal ageing on the tensile and impact properties of ODS–18Cr ferritic steel. Given the very limited work carried out till date on the mechanical behavior of ODS–18Cr steel, there is a clear need to understand their deformation behavior at different temperatures and also to model their mechanical behavior.
Fig.1. Fe–Cr phase diagram [1].
3
2.
Experimental Details
Inert gas atomized pre–alloyed 18Cr ferritic steel (Fe–18Cr–2.3W–0.3Ti) powder with an average particle size of 100 µm (produced at International Advanced Research Centre for Powder Metallurgy and New Materials, Hyderabad, India) was milled with and without nano yttria powder (M/s. Inframat Advanced Materials, USA, size: 30–50 nm, purity: 99.95 %) in a high energy horizontal attritor mill (Simoloyer CM–20) up to 6 h. Milling was carried out at room temperature in a water cooled stainless steel container using hardened steel balls of 5 mm under argon atmosphere. Ball to powder ratio of 7.5:1 and milling speed of 550 rpm were maintained for all milling experiments. The milled powders were filled in mild steel cans, degassed and vacuum sealed. The sealed cans were upset forged at 1323 K and hot extruded at 1423 K with an extrusion ratio of 19. The extruded rods were annealed at 1173 K for 1 h and water quenched. The finer details of the process were given in the earlier publication of the authors [16].
The chemical composition prealloyed powder as well as bulk samples was evaluated using ICPAES (JOBIN YUON FRANCE, Model: Ultima2CHR), LECO oxygen/nitrogen (Model: TC436) and LECO carbon/sulpher (Model: CS444) analyzers. The chemical compositions of pre–alloyed 18Cr steel powder, extruded and annealed oxide free 18Cr steel (NODS) and ODS18Cr steel are given in Table 1. Grain sizes of all the annealed rods were estimated from the grain boundary maps obtained by Electron Back–Scattered Diffraction (EBSD) technique available with FESEM (Hitachi, model: S4300SE/N). Grains with misorientation angle >15 are only considered for the calculation of the average size. Transmission electron microscopic (TEM) observations of annealed samples before and after tensile testing were performed using FEI Tecnai G2 200 kV (LaB6) microscope equipped with Olympus (SIS mega view III) wide angle camera. Dispersoid size, number density and 4
volume fraction calculations were carried out as per the procedures reported by the authors in their earlier publication [16].
The tensile properties of the annealed samples were evaluated as per ASTM E21 standard from room temperature (RT (298 K)) to 1073 K at a strain rate of 7.5×10–4 s–1 using Universal Testing Machine (INSTRON, Model: 4507, 200 KN). Round tensile test specimens of 3.99 mm and 22 mm gauge length were used. A constant heating rate of 10 C/min, hold time of 2 h at test temperature and furnace cooling was maintained for tensile testing of all the samples at temperatures above 298 K. For each test condition, the average results obtained on two specimens are reported. Sample/ Composition Pre–alloyed 18Cr steel powder
Fe
Cr
W
Ti
C
Total O
N
Y2O3
Excess O
bal.
17.8
2.3
0.34
<0.03
0.006
0.003
–
–
NODS
bal.
17.6
2.3
0.34
<0.03
0.05
0.012
–
–
ODS–F
bal.
17.4
2.2
0.31
<0.03
0.14
0.012
0.36
0.06
Table 1 Chemical composition (wt. %) of pre–alloyed powder, NODS and ODS–F samples.
3.
Results
3.1.
Chemical Composition
The chemical analysis indicated that the oxygen and nitrogen contents in both ODS and NODS steels increased slightly when compared to the pre–alloyed steel powder due to minor leakages during milling in argon atmosphere. Since the Y/ Ti atomic ratio continues to be 1:1 in spite of the O pick up, its importance to the overall structure and properties is negligible [33].
5
3.2.
Microstructural Features
The evolution of microstructure in ODS–18Cr steels during various processing steps (milling, upset forging, extrusion and annealing) has been studied exhaustively by the authors and published elsewhere [16]. However, for the sake of completeness, the essential microstructural features in extruded and annealed condition as a function of milling time are presented. Typical variation of TEM microstructures of ODS–18Cr steel milled for 1, 3 and 6 h of milling and of NODS obtained on the longitudinal sections is shown in Figs. 2 (a–d) respectively. The grains are elongated at 1 h milling but are more equiaxed in the case of 3 and 6 h milled ODS samples. NODS (Fig. 2 (a)) exhibits very coarse and uniform grain size. The average grain sizes (longitudinal direction) of extruded and annealed NODS and ODS steels estimated by EBSD are given in Table 2. Additionally, TEM bright field images of ODS steels milled for 1, 3 and 6 h which reveal the dispersoids are presented in Figs. 3 (a–c) respectively. The average size, number density and volume fraction of dispersoids estimated by detailed TEM examination are also included in Table 2.
6
Fig. 2. TEM images showing the grain structure of extruded+annealed samples in longitudinal direction: (a) ODS–C, (b) ODS–M, (c) ODS–F and (d) NODS.
Name
Average grain size (dg), [µm] (longitudinal)
Average dispersoid size (dp)[nm]
Number density [m–3]
Volume fraction
Inter– dispersoid spacing (S), [nm]
NODS
13 (± 6.28)
–
–
–
–
ODS–C
1.17 (± 0.27)
15 (± 11.76)
6.21x1021
0.0092
101
ODS–M
0.73 (± 0.26)
10 (± 7.03)
1.82x1022
0.0094
66
5 (± 2.35)
1.24x1023
0.0096
33
0.48 (± 0.21) 0.30 (± 0.21)* * In transverse direction ODS–F
Table 2 Average grain size, average size, number density, volume fraction and inter– dispersoid spacing of dispersoids in extruded+annealed ODS–18Cr steels.
7
Fig. 3. TEM bright field images of extruded+annealed ODS–18Cr steels: (a) ODS–C, (b) ODS–M and (c) ODS–F.
In Fig. 4, the relationship between milling times and grain size and dispersoid size using the data in Table 2 is presented in a graphical form. This figure clearly shows that increasing milling time (1 to 6 h) decreases both grain size and dispersoid size simultaneously. Henceforth, ODS–18 Cr steel milled for 1, 3 and 6 h will be designated as ODS–C, ODS–M and ODS–F with the letters C, M and F denoting coarse, medium and fine microstructures.
8
Fig. 4. The contribution of grain size and dispersed size obtained for 1, 3 and 6 h milling.
To assess the combined effect of strain and temperature on the extent of coarsening of the grains and dispersoids during the tensile test, microstructural examination was carried out on ODS–F sample after tensile deformation at 298 and 1073 K. The requisite sample in transverse direction was extracted from around 1 mm below the fracture surface for the above purpose. The relevant microstructural features showing the grain structures and dispersoids of ODS–F in annealed and deformed (298 and 1073 K) conditions are presented in Figs. 5 (a– c)–(d–f) respectively. The corresponding quantified data of grain and dispersoid sizes of ODS–F sample are given in Table 3. The data presented clearly indicates that neither the grains nor the dispersoids exhibit any coarsening even when exposed to a temperature of 1073 K during the tensile test. It is well established that the dispersoids are formed by the precipitation of Y–Ti–O complexes during hot consolidation either from the ferrite containing Y, Ti and O in solid solution or from nano clusters in as milled condition [34]. All 9
the solutes are expended in the formation of stable Y–Ti–O complex oxides during high temperature exposure, leaving little or no Y left in solid solution to promote growth of dispersoids. Further, because of the stability of Y–Ti–O complex oxide dispersoids, coarsening does not occur up to 1200 K even if assisted by strain. This observation implies that the grain size, dispersoid size and their number density can be assumed to be invariant with test temperature while modeling the yield strength behavior of ODS–C, ODS–M and ODS–F steels.
Fig. 5. TEM images showing the grain structures and dispersoids of ODS–F in (a and d annealed condition, (b and e) deformed at RT (298 K) and (c and f) deformed at 1073 K, respectively.
Sample name
ODS–F
Sample condition
Average Average grain size (dg), [µm] dispersoid Size (dp) (transverse direction) [nm]
Annealed
0.32 (± 0.18)
4.9 (± 2.35)
Tensile tested at 298 K
0.3 (± 0.14)
4.5 (± 1.79)
Tensile tested at 1073 K
0.31 (± 0.17)
4.6 (± 1.93)
Table 3 Average grain size (transverse) and dispersoid size in ODS–F in annealed sample and in deformed condition at RT (298 K) and 1073 K.
10
3.3.
Tensile Properties
Tensile tests were carried out on ODS–18Cr and NODS steels from 298 to 1073 K to assess the effect of microstructure, resulting from varying milling times, on strength and ductility. The engineering stress–strain curves of NODS and ODS–18Cr steels at different test temperatures are shown in Figs. 6 (a–d) respectively. The variation of yield strength (y) and ultimate tensile strengths (u) of ODS–C, ODS–M and ODS–F along with NODS with temperature is shown in Figs. 7 (a–b) respectively. The uniform (eu) and total fracture strains (et) for NODS and ODS–18Cr steels are shown in Figs. 8 (a–b) respectively. From the tensile data presented in Figs. 6 to 8, the following conclusions can be drawn. (a) Fig. 6 indicates that at low test temperatures (673 K and below), the shape of the tensile curves of both ODS–18Cr and NODS steels changed from characteristic yielding to work hardening to softening (necking) and finally fracture. At temperatures above 673 K, the work hardening phase was absent. (b) The engineering stress–strain curves (Fig. 6) of all the alloys (including NODS) indicate that the decrease in yield strength and ultimate tensile strength (UTS) with increasing temperatures is marginal up to 673 K but more dramatic at higher temperatures. Such behavior is clearly seen in Figs. 7 as well. (c) The effect of progressive refining of the microstructure (i.e, ODS–C to ODS–M to ODS–F) is to increase the yield strength and UTS at all temperatures–even at 800 C (1073 K). Such behavior can be attributed to the thermal stability of the grains and dispersoids even at 1073 K as observed in Fig. 5. (d) The ductility values of ODS alloys (Fig. 8 (a)), covering the temperature range 298 to 1073 K, lie in the range 20 to 35 % which is quite acceptable. NODS alloy exhibits higher ductility as compared to ODS alloys. The difference in fracture strain between
11
NODS and ODS alloys is due to the drastic reduction in strain hardening in ODS alloys which in turn is caused by the fine grained structure. The unusually large fracture strain in NODS at above 873 K is attributable to near zero strain hardening exponent and favourable temperature, grain size and strain rate hardening. The ODS– 18Cr steels on the other hand showed decrease in fracture strain due to the stability of microstructure which promotes inhomogeneous deformation and pinning of grain boundaries by oxide particles inhibiting grain boundary sliding even at high temperatures. (e) In all alloys (including NODS), the uniform strain decreases progressively with increasing test temperature (see Fig. 8 (b)). NODS exhibit the highest uniform strain values and ODS–F the lowest.
Fig. 6. Engineering stress–strain curves of 18Cr and ODS–18Cr steels at different test temperatures: (a) NODS, (b) ODS–C, (c) ODS–M and (d) ODS–F.
12
Fig. 7. Variation of (a) y and (b) u of 18Cr and ODS–18Cr steels with test temperature.
Fig. 8. Variation of ductility indices of (a) eu and (b) et of 18Cr and ODS–18Cr steels with test temperature.
4.
Discussion
4.1.
Microstructural Changes with Milling Time
An unique feature of the present work is the demonstration that using milling time as one of the process variables, one can obtain a range of microstructures (coarse to medium to fine) in the same ODS–18Cr alloy. The events leading to the refinement of microstructure with increasing milling time are reasonably well understood. Many studies [16, 35–37] have suggested that, Y2O3 particles added to the pre–alloyed 18Cr steel powder fragment progressively with increasing milling time and ultimately dissolve in the matrix. Simultaneously, nano clusters of Y–Ti–O start forming during the final stages of milling. These Y, Ti and O nano clusters transform to Y–Ti–O nano dispersoids during subsequent 13
hot extrusion and annealing. Higher the milling time, finer and more numerous are the nanoclusters and hence finer are the dispersoids even after extrusion and annealing. Thus, the effect of milling time persists on the microstructure even after subsequent extrusion and annealing. The sequence of events with regard to grain size is similar. With increasing milling time, the grain size reduces in size but some grain growth occurs during subsequent hot extrusion and annealing. However, the influence of milling time still persists. It is also clear that the nano dispersoids are very effective in pinning the grain boundaries resulting in extremely fine grain size in ODS–18Cr steel (0.48 to 1.17 m) as compared to NODS wherein the final grain size is an order of magnitude higher at 13 m.
4.2.
Influence of Microstructure and Temperature on Yield Strength
The strong influence of microstructure on yield strength at all temperatures (298 to 1073 K) is evident from Table 4, wherein the yield strength of ODS–C, ODS–M and ODS–F is presented at various temperatures. The increase in yield strength due to refinement of microstructure (σy(F)–σy(C), Table 4) is 351 MPa at 298 K and decreases with increasing temperature to 116 MPa at 1073 K. Looked another way, the ratio of yield strength of ODS–F to ODS–C (see Table 4) is around 1.76 at room temperature (298 K) and increases to 2.0 at 1073 K. In contrast, the ratio of yield strength at 1073 to 298 K, in the case of ODS–C, ODS– M and ODS–F alloys, provided in the bottom row of Table 4, is reasonably constant (0.25 to 0.28). Thus, the influence of microstructure persists even at the highest test temperature of 1073 K. All these aspects of microstructure–yield strength correlation will be understood on the basis of a model for yield strength proposed in the next section.
14
σy [MPa] ∆σy [MPa] σy (ODS–F) Temperature [K] C M F σy(F)–σy(C) σy (ODS–C) 298 459 701 810 351 1.76 673 377 549 662 285 1.76 873 244 375 453 209 1.86 1073 116 192 232 116 2.00 σy (1073 K) 0.25 0.27 0.28 – – σy (298 K) Table 4 The yield strength of ODS–18Cr steels as a function of temperature and microstructure (C, M and F).
4.3.
Modeling of the Yield Strength at Room and Elevated Temperatures
The possible strengthening mechanisms contributing to the yield strength of ODS–18Cr steels include Peierls–Nabarro lattice resistance stress, solid solution strengthening, dislocation forest hardening, Hall–Petch grain boundary strengthening and dispersion strengthening due to Orowan mechanism.
In the present study, the yield strength of not only the ODS–18Cr steels (C, M and F) but also of NODS steel, which does not have the dispersoids, have been experimentally obtained. Thus, the difference in yield stress between ODS–18Cr steel and NODS steel is largely due to Hall–Petch grain boundary strengthening and dispersion strengthening due to Orowan mechanism. The extent of strengthening due to Peierls–Nabarro lattice resistance, solid solution strengthening and forest dislocation hardening should be the same in both NODS and ODS–18Cr steel. This is especially true because the solid solution strengthening elements like W, Cr do not get incorporated in the dispersoid and thus are available for solid solution strengthening to the same extent in NODS and ODS–18Cr steels. Additionally, since both NODS and ODS–18Cr steel have been annealed at 1173 K after extrusion, it is likely that the dislocation density exhibits low and similar values in both materials. However, NODS also includes some contribution from Hall–Petch strengthening since it has an average grain size 15
of 13 m. Thus, the matrix strength (m), which includes only the contributions from Peierls lattice stress, solid solution strengthening and forest dislocation hardening is now given as,
m y ,NODS H P,NODS
(1)
Finally, the yield strength of ODS–18Cr steels is obtained as,
y ,ODS m H P Or
(2)
In Eq. (2), ∆σH–P represents the yield strength increment due to grain boundary strengthening (Eq. (3)) and ∆σOr is the yield strength increment due to dispersion strengthening (Eqs. (4) and (5)). H P K H P d g1/ 2 ,
(3)
Gb d p Or A ln , S 2b
(4)
S
2 3 f 2
dp
1/ 2
22 3
1/ 2
,
(5)
To calculate the yield strength at all test temperatures (298 to 1073 K), it is important to understand the influence of temperature on the various parameters used in Eqs. (3)–(5). We have already demonstrated that grain size and dispersoid size does not vary even up to 1073 K (see Table 3 and Fig. 5). Thus, the values of dg, dp, S and f listed in Table 5 are valid at all temperatures. The shear modulus of the alloys decreases with increasing temperature as per Eq. (6) given below [38].
G(GPa) 93.2(1 4.368 104 T )
(6)
Another parameter which changes with temperature is the Hall–Petch constant (KH–P) and for the present alloy, the data of Hall–Petch constant vs. temperature provided by Kim et al. [38] in the case of ODS–14Cr steel has been utilized. Lastly, the matrix strength (σm) also varies with temperature. However, the yield strength of NODS alloy as a function of temperature
16
has been obtained in the present study and this data directly can be utilized for determining σm (as per Eq. (1)) at various temperatures.
Parameter
Values chosen NODS
ODS–C
ODS–M
ODS–F
Ref.
A
Numerical constant
–
0.3
0.3
0.3
45
b
Burgers vector [nm]
0.25
0.25
0.25
0.25
45
KH–P
Hall–Petch constant [MPa.m1/2]
338
338
338
338
38
dg
Grain size [µm]
13.0
1.17
0.73
0.48
Present study
dp
Dispersoid size [nm]
–
15
10
5
Present study
S
Inter–dispersoid spacing [nm]
–
101
66
33
Present study
f
Volume fraction of dispersoids
–
0.0092
0.0094
0.0096
Present study
Table 5 The values chosen for the various parameters appearing in the model (Eqs. (1)–(5)) for NODS and ODS–18Cr steels and the reference source for the data.
On the basis of the equations and values for the various parameters detailed above, it is now possible to calculate the variation of yield strength of ODS–Cr steels (σy,ODS; Eq. (2)) with temperature and as a function of microstructure (C, M and F). At this stage it is important to note that Eq. (2) used to calculate σy,ODS assumes that the yield stress increments due to various strengthening mechanisms can be linearly summed up and this in turn implies that there is no influence of one strengthening mechanism on the other. The prediction of the linear summation model as a function of both temperature and microstructure is presented in Fig. 9. It is clear that linear summation predicts yield strength values which are considerably higher than the experimental values at all test temperatures and at different microstructures. This is not surprising since dislocations participate in overcoming many of the obstacles present in the ODS alloys and availability of dislocations for participating in any one
17
strengthening mechanism is limited and dependent on other strengthening mechanisms which are also operating at the same time.
A number of researchers have analyzed the problem of summation of various strengthening mechanisms [39–47]. Based on extensive simulation of dislocation motion in the presence of multiple obstacles, Brown and Ham [39] have concluded that linear summation is a poor approximation except when a few strong obstacles are dispersed among a large number of weak obstacles and that quadratic summation has more validity. A generalized version of the non–linear summation is given by Eq. (7) given below.
yk 1k 2k 3k
(7)
In Eq. (7), σ1, σ2 and σ3 represent 3 different strengthening mechanisms. The value of parameter ‘k’ defines the nature of interaction between the 3 strengthening mechanisms. A value of 1 for k, transforms Eq. (7) to one representing linear summation like Eq. (2). On the other hand, k=2 represents root mean square or quadratic summation and implies that the strengthening mechanisms have synergistic interaction.
Koppenaal and Kulhman–Wilsdorf [40] have proposed the quadratic summation (k=2) model which has been supported by subsequent statistical calculations and simulations. Simulation of combined effect of Orowan and forest hardening by Queyreau et al. [41] indicated a value of 2 for k as most suitable. Lagerpusch et al. [42] investigated experimentally the superposition of solid solution strengthening and dispersion strengthening in a Cu–Au–SiO2 alloy and concluded that k=1.8 gives the best fit of the experimental data. It has been shown that when the obstacle strengths and number density are comparable, a value of 2 for k is the most appropriate [43]. Hansen [44] has demonstrated that superposition of grain boundary and particle strengthening is best explained by assuming a value of 2 for k. Nembach [45] in 18
an interesting study of the interaction between grain boundaries and precipitates in a nickel base superalloy, observed that when the precipitate size is less than 20 nm, most of the dislocations were piled up at grain boundaries and only a few Orowan loops were observed. In contrast, when the precipitate size is greater than 25 nm, only Orowan loops were seen. Thus, the interaction between two strengthening mechanisms can be quite complex.
Given the above scenario and also since the obstacles in the present ODS–18Cr alloys (grain boundary and dispersoids) have comparable strength, a k value of 2 is the most appropriate. Thus, the equation for yield strength of ODS alloy can be written down as, 2 y ,ODS ( rmss) ( m2 H2 P Or )
1
2
(8)
In the above equation, the yield strength is referred to as σy,ODS (rmss) wherein the term rmss stands for root mean square superposition.
A few investigators [46, 47] have suggested that the resistance experienced by dislocations to Peierls lattice stress, solid solution and forest dislocations can be treated as a friction stress impeding dislocation motion and in such a case linear superposition (k=1) is appropriate for some of the strengthening mechanisms. The above aspect can be taken into account in the present alloy by adding matrix strength (σm) to the root mean square of grain boundary and dispersion strengthening as indicated below. 2 y ,ODS ( prmss) m ( H2 P Or )
1
2
(9)
In Eq. (9), the yield strength is referred to as σy,ODS (prmss) wherein the term prmss stands for part root mean square summation.
19
Finally, the variation of the calculated yield strength with temperature and microstructure assuming rms summation and part rms summation are presented in Fig. 9. This figure clearly suggests that rms superposition predicts the yield strength most accurately as compared to the
predictions assuming part rms or linear superposition. Fig. 9. Comparison of the experimentally obtained yield strength as a function of temperature of ODS–18Cr steel for three microstructural conditions (ODS–C, ODS–M & ODS–F) with the predictions of yield strength assuming linear, root mean square (rms) and partial rms superposition.
Now that the ability of the model to predict the yield strength as a function of temperature and microstructure has been validated, it is possible to assess the relative contributions of grain boundary strengthening (∆σH–P) and dispersion strengthening (∆σOr). The ratio of ∆σH–P to ∆σOr is presented in Fig. 10 as a function of temperature and for three microstructures (ODS–C, ODS–M and ODS–F). This figure clearly indicates that grain boundary strengthening contribution to yield strength is greater than that of dispersion strengthening up 20
to a temperature of 873 K. At the highest temperature of 1073 K the contribution of dispersion strengthening is substantially higher than grain boundary strengthening. This can be explained on the basis of weakening of the grain boundaries at temperature higher than 50 % of the melting point. In contrast, the variation of dispersion strengthening follows the temperature dependence of the shear modulus since the dispersoids are stable even at this temperature.
Fig. 10. The variation of the ratio ∆σH–P / ∆σOr with temperature for ODS–18Cr with coarse (C), medium (M) and fine (F) microstructure.
5.
Conclusions
The tensile deformation behavior of ODS–18Cr ferritic steels along with 18Cr (NODS) steel at indifferent test temperatures have been investigated. The salient features from the present study are given below: The strength of ODS–18Cr steels increased at all the test temperatures due to the refinement of dispersoids and grains. 21
All ODS–18Cr steels exhibited superior strength values over the whole range of test temperatures compared to NODS steel and the strength of both the steels decreased with temperature. The ductility indices of ODS–18Cr steels showed marginal decrease with refinement of microstructure over the test temperature range and were lower than those of NODS steel. With increasing test temperature, NODS and ODS–18Cr steels exhibited continuous decrease in uniform elongation, but the fracture strain remained nearly same. The strength values, predicted by root mean square superposition model at room and higher temperatures compare well with experimental values.
Acknowledgements The authors thank Dr. G. Ravi Chandra, Mr. M. Ramakrishna and Mr. G.V.R. Reddy for the help for EBSD, TEM and SEM studies. They also thank Dr. K. Satya Prasad for his help in carrying out microstructural examination. The authors express their gratitude to IGCAR, Kalpakkam for funding (No. IGC/MMG/MMD/ODS/01/2010) the work and NFC, Hyderabad for carrying out hot extrusion.
References [1]
ASM Handbook Vol. 3: Alloy Phase Diagrams, ASM International, Materials Park, Ohio, 1998, 44073–0002.
[2]
G.R. Odette, M.J. Alinger, B.D Wirth, Recent developments in irradiation–resistant steels, Annu. Rev. Mater. Res. 38 (2008) 471–503.
[3]
S. Ukai, M. Fujiwara, Perspective of ODS alloys application in nuclear environments, J. Nucl. Mater. 307–311 (2002) 749–757.
[4]
A. Ramar, P. Spatig, R. Schaublin, Analysis of high temperature deformation mechanism in ODS EUROFER97 alloy, J. Nucl. Mater. 382 (2008) 210–216. 22
[5]
S. Ohtsuka, S. Ukai, M. Fujiwara, T. Kaito, T. Narita, Improvement of creep strength of 9CrODS martensitic steel by controlling excess oxygen and titanium concentrations, Mater. Trans. 46 (2005) 487–492.
[6]
R.L. Klueh, J.P. Shingledecker, R.W. Swindeman, D.T. Hoelzer, Oxide dispersion– strengthened steels: A comparison of some commercial and experimental alloys, J. Nucl. Mater. 341 (2005) 103–114.
[7]
S. Ukai, Oxide dispersion strengthened steels, Comp. Nucl. Mater. 4 (2012) 241–271.
[8]
S. Ukai, S. Mizuta, T. Yoshitake, T. Okuda, M. Fujiwara, S. Hagi, T. Kobayashi, Tube manufacturing and characterization of oxide dispersion strengthened ferritic steels, J. Nucl. Mater. 283–287 (2000) 702–706.
[9]
A. Alamo, V. Lambard, X. Averty, M.H. Mathon, Assessment of ODS–14%Cr ferritic alloy for high temperature applications, J. Nucl. Mater. 329–333 (2004) 333–337.
[10]
T. Jayakumar, M.D. Mathew, K. Laha, High temperature materials for nuclear fast fission and fusion reactors and advanced fossil power plants, Procedia Eng. 55 (2013) 259–270.
[11]
H.S. Cho, A. Kimura, S. Ukai, M. Fujiwara, Corrosion properties of oxide dispersion strengthened steels in super–critical water environment, J. Nucl. Mater. 329–333 (2004) 387–391.
[12]
A. Kimura, H.S. Cho, N. Toda, R. Kasada, K. Yutani, H. Kishimoto, N. Iwata, S. Ukai, M. Fujiwara, High burn up fuel cladding materials R&D for advanced nuclear systems: Nano–sized oxide dispersion strengthening steels, J. Nucl. Sci. Technol. 44 (2007) 323–328.
[13]
T. Kaito, T. Narita, S. Ukai, Y. Matsuda, High temperature oxidation behavior of ODS steels, J. Nucl. Mater. 329–333 (2004) 388–1392.
[14]
J.H. Kim, K.M. Kim, T.S Byun, D.W. Lee, C.H. Park, High temperature oxidation behavior of nano–structured ferritic oxide dispersion strengthened alloys, Thermochem. Acta 579 (2014) 1–8.
[15]
C. Capdevila, M. Serrano, M. Campos, High strength oxide dispersion strengthened steels: fundamentals and applications, Mater. Sci. Technol. 30 (2014) 1655–1657.
[16]
M. Nagini, R. Vijay, K.V. Ralulapati, K. Bhanu Sankara Rao, M. Ramakrishna, A.V. Reddy, G. Sundararajan, Effect of process parameters on microstructure and hardness of oxide dispersion strengthened 18Cr ferritic steel, Metall. Mater. Trans. A, 47A (2016) 4197–4209.
[17]
J.H. Lee, Development of oxide dispersion strengthened ferritic steels with and without aluminum, Front. Energy 6 (2012) 29–34.
[18]
Y. de Carlan, J.L. Bechade, P. Dubuisson, J.L Seran, P. Billot, A. Bougault, T. Cozzika, S. Doriot, D. Hamon, J. Henry, M. Ratti, N. Lochet, D. Nunes, P. Olier, T. Leblond, M.H. Mathon, CEA developments of new ferritic ODS alloys for nuclear applications, , J. Nucl. Mater. 386–388 (2009) 430–432. 23
[19]
H. Hadraba, B. Fournier, L. Stratil, J. Malaplate, A.L. Rouffie, P. Wident, L. Ziolek, J.L. Bechade, Influence of microstructure on impact properties of 9–18%Cr ODS steels for fusion/fission applications, J. Nucl. Mater. 411 (2011) 112–118.
[20]
Deepak Kumar, Ujjwal Prakash, Vikram V. Dabhade, K. Laha, T. Sakthivel, Development of oxide dispersion strengthened (ODS) ferritic steel through powder forging, J. Mater. Eng. Perform. 26 (2017) 1817–1824.
[21]
A.L. Rouffie, J. Crepin, M. Sennour, B. Tanguy, A. Pineau, D. Hamon, P. Wident, S. Vincent, V. Garat, B. Fournier, Effect of the thermal ageing on the tensile and impact properties of a 18% Cr ferritic steel, J. Nucl. Mater. 445 (2014) 37–42.
[22]
S. Li, Z. Zhou, M. Li, M. Wang, G. Zhang, Microstructure characterization and tensile properties of 18Cr–4Al oxide dispersion strengthened ferritic steel, J. Alloys Comp. 648 (2015) 39–45.
[23]
P. Dubuisson, Y. de Carlan, V. Garat, M. Blat, ODS Ferritic/martensitic alloys for sodium fast reactor fuel pin cladding, J. Nucl. Mater. 428 (2012) 6–12.
[24]
S. Li, Z. Zhou, J. Jang, M. Wang, H. Hu, H. Sun, L. Zou, G. Zhang, L. Zhang, The influence of Cr content on the mechanical properties of ODS ferritic steels, J. Nucl. Mater. 455 (2014) 194–200.
[25]
B. Fournier, A. Steckmeyer, A.L. Rouffie, J. Malaplate, J. Garnier, M. Ratti, P. Wident, L. Ziolek, I. Tournie, V. Rabeau, J.M. Gentzbittel, T. Kruml, I. Kubena, Mechanical behaviour of ferritic ODS steels–Temperature dependency and anisotropy, J. Nucl. Mater. 430 (2012) 142–149.
[26]
J.H. Kim, T.S. Byun, J.H. Lee, J.Y. Min, S.W. Kim, C.H. Park, B.H. Lee, Effects of processing condition on the microstructural and tensile properties of 14Cr–based oxide dispersion strengthened alloys, J. Nucl. Mater. 449 (2014) 300–307.
[27]
P. He, R. Lindau, A. Moeslang, H.R.Z. Sandim, The influence of thermomechanical processing on the microstructure and mechanical properties of 13.5Cr ODS steels, Fusion Eng. Des. 88 (2013) 2448–2452.
[28]
R. Didomizio, S. Huang, L. Dial, J. Ilavsky, M. Larsen, An assessment of milling time on the structure and properties of a nanostructured ferritic alloy (NFA), Metall. Mater. Trans. A, 45 A (2014) 5409–5418.
[29]
A. Garcia–Junceda, M. Hernandez–Mayoral, M. Serrano, Influence of the microstructure on the tensile and impact properties of a 14Cr ODS steel bar, Mater. Sci. Eng. A, 556 (2012) 696–703.
[30]
M. Serrano, M.H. Mayoral, A.G. Junceda, Microstructural anisotropy effect on the mechanical properties of a 14Cr ODS steel, J. Nucl. Mater. 428 (2012) 103–109.
[31]
A. Steckmeyer, M. Praud, B. Fournier, J. Malaplate, J. Garnier, J.L. Bechade, I. Tournie, A. Tancray, A. Bougault, P. Bonnaillie, Tensile properties and deformation mechanisms of a 14Cr ODS ferritic steel, J. Nucl. Mater. 405 (2010) 95–100.
24
[32]
M. Praud, F. Mompiou, J. Malaplate, D. Caillard, J. Garnier, A. Steckmeyer, B. Fournier, Study of the deformation mechanisms in a Fe–14% Cr ODS alloy, J. Nucl. Mater. 428 (2012) 90–97.
[33]
S. Ukai, S. Mizuta, M. Fujiwara, T. Okuda, T. Kobyashi, Development of 9Cr–ODS martensitic steel claddings for fuel pins by means of ferrite to austenite phase transformation, J. Nucl. Sci. Technol. 39 (2002) 778–788.
[34]
P. Miao, G.R. Odette, T. Yamamoto, M. Alinger, D. Klingensmith, Thermal stability of nano–structured ferritic alloy, J. Nucl. Mater. 377 (2008) 59–64.
[35]
M. Ratti, D. Leuvrey, M.H. Mathon, Y. de Carlan, Influence of titanium on nano– cluster (Y, Ti, O) stability in ODS ferritic materials, J. Nucl. Mater. 386–388 (2009) 540–543.
[36]
C. Cayron, A. Montani, D. Venet, Y. de Carlan, Identification of new phases in annealed Fe–18CrWTi ODS powders, J. Nucl. Mater. 399 (2010) 219–224.
[37]
Olena Kalokhtina, B. Radiguet, Y. de Carlan, P. Pareige, Study of nano–cluster formation in Fe–18Cr ODS ferritic steel by atom probe tomography, Mater. Res. Soc. Symp. Proc. 1264 (2010) 1264–BB03–04.
[38]
J.H. Kim, T.S Byun, D.T. Hoelzer, C.H. Park, J.T. Yeom, J.K. Hong, Temperature dependence of strengthening mechanisms in the nanostructured ferritic alloy 14YWT: Part II–Mechanistic models and predictions, Mater. Sci. Eng. A, 559 (2013) 111–118.
[39]
L.M. Brown, R.K. Ham, Dislocation–particle interactions, in: A. Kelly, R.B. Nicholson (Eds), Strengthening methods in solids, Applied Science Publishers, London, 1971, 9–135.
[40]
T.J. Koppenaal, D. Kuhlmann–Wilsdorf, The effect of prestressing on the strength of neutron–irradiated copper single crystals, Appl. Phys. Lett. 4 (1964) 59–61.
[41]
S. Queyreau, G. Monnet, B. Devincre, Orowan strengthening and forest hardening superposition examined by dislocation dynamics simulations, Acta Mater. 58 (2010) 5586–5595.
[42]
U. Lagerpusch, V. Mohles, E. Nembach, On the additivity of solid solution and dispersion strengthening, Mater. Sci. Eng. A 319–321 (2001) 176–178.
[43]
Y. Dong, T. Nogaret, W.A. Curtin, Scaling of dislocation strengthening by multiple obstacle types, Metall. Mater. Trans. A 41 A (2010) 1954–1960.
[44]
N. Hansen, Super position of Grain–boundary strengthening and particle strengthening in aluminium, Scr. Mater. 5 (1971) 417–420.
[45]
E. Nembach, Synergistic effects in the super position of strengthening mechanisms, Acta Mater. 40 (1992) 3325–3330.
[46]
M. Dade, J. Malaplate, J. Garnier, F. De Geuser, F. Barcelo, P.Wident, A. Deschamps, Acta Materialia, Influence of microstructural parameters on the
25
mechanical properties of oxide dispersion strengthened Fe–14Cr steels, Acta Mater. 127 (2017) 165–177. [47]
J. Lu, O. Omotoso, J.B. Wiskel, D.G. Ivery, Strengthening mechanisms and their relative contributions to the yield strength of micro alloyed steels, Metall. Mater. Trans. A 43 A (2012) 3043–3061.
26
TABLES Table 1.
Chemical composition (wt. %) of pre–alloyed powder, NODS and ODS–F samples.
Table 2.
Average grain size, average size, number density, volume fraction and inter– dispersoid spacing of dispersoids in extruded+annealed ODS–18Cr steels.
Table 3.
Average grain size (transverse) and dispersoid size in ODS–F in annealed sample and in deformed condition at RT (298 K) and 1073 K.
Table 4.
The yield strength of ODS–18Cr steels as a function of temperature and microstructure (C, M and F).
Table 5.
The values chosen for the various parameters appearing in the model (Eqs. (1)–(5)) for NODS and ODS–18Cr steels and the reference source for the data.
FIGURE CAPTIONS Fig. 1. Fe-Cr phase diagram. Fig. 2. TEM images showing the grain structure of extruded+annealed samples in longitudinal direction: (a) ODS–C, (b) ODS–M, (c) ODS–F and (d) NODS. Fig. 3. TEM bright field images of extruded+annealed ODS–18Cr steels: (a) ODS–C, (b) ODS–M, (c) ODS–F Fig. 4. The contribution of grain size and dispersoid size obtained for 1, 3 and 6 h milling. Fig. 5. TEM images showing the grain structures and dispersoids of ODS–F in (a and d) annealed condition, (b and e) deformed at RT (298 K) and (c and f) deformed at 1073 K, respectively. Fig. 6. Engineering stress–strain curves of 18Cr and ODS–18Cr steels at different test temperatures: (a) NODS, (b) ODS–C, (c) ODS–M and (d) ODS–F. Fig. 7. Variation of (a) y and (b) u of 18Cr and ODS–18Cr steels with test temperature. Fig. 8. Variation of ductility indices of (a) eu and (b) et of 18Cr and ODS–18Cr steels with test temperature. Fig. 9. Comparison of the experimentally obtained yield strength as a function of temperature of ODS–18Cr steel for three microstructural conditions (ODS–C, ODS– M and ODS–F) with the predictions of yield strength assuming linear, root mean square (rms) and partial rms superposition. Fig.10. The variation of the ratio ∆σH–P /∆σOr with temperature for ODS–18Cr steel with coarse (C), medium (M) and fine (F) microstructure. 27
Graphical abstract
High lights:
Tensile deformation behavior of nano oxide dispersoid strengthened 18Cr ferritic steels was studied.
The grain size of ODS–18Cr steels progressively decreased with the size of dispersoids.
Presence of nano sized dispersoids enhanced the strength at all temperatures when compared to dispersoid free 18Cr steel.
The flow strength of ODS–18Cr steels also increased with refinement of microstructure (grain and dispersoid size) at all test temperature.
Microstructure–mechanical property correlations were developed using root mean square superposition model.
28