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Phase transformations and mechanical properties in heat treated superaustenitic stainless steels
Materials Science & Engineering A 561 (2013) 477–485
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Phase transformations and mechanical properties in heat treated superaustenitic stainless steels T. Koutsoukis a,n, A. Redjaı¨mia b, G. Fourlaris a a b
Laboratory of Physical Metallurgy, National Technical University of Athens, 9 Heroon Polytechniou Street, 15780 Athens, Greece Universite´ de Lorraine, Institut Jean Lamour, UMR-CNRS 7198, Nancy F-54042, France
a r t i c l e i n f o
abstract
Article history: Received 27 June 2012 Received in revised form 16 October 2012 Accepted 17 October 2012 Available online 26 October 2012
A microstructure–properties relationship study in two superaustenitic stainless steels (S31254 and S32654) was carried out, following exposure at elevated temperatures for various ageing times. Due to high temperature ageing, most stainless steel grades suffer the formation of various precipitates, directly affecting their properties. The full characterization of those precipitates and the correlation with the mechanical behavior of the steels is the primary aim of this study. Samples of the steel grades studied, were exposed to isothermal heat treatments within the temperature range of 650–950 1C, for ageing times varying between 0.5 h and 3000 h, followed by water quenching at room temperature. Microstructural examination indicated the formation of four different secondary phases, sigma phase (s), chi phase (w), Laves phase and b-Cr2N nitride, which were characterized by transmission electron microscopy (TEM) and electron diffraction. The results obtained permitted the construction of the time–temperature–precipitation (TTP) plots. In addition, tensile and Vickers hardness testing were utilized and the modulus of toughness was calculated. The kinetics of the formation of various precipitates with increasing temperature and aging duration was also observed. It was found that various precipitates had a significant effect on all mechanical properties studied. & 2012 Elsevier B.V. All rights reserved.
Keywords: Superaustenitic stainless steel Ageing Precipitation Laves phase Mechanical properties TEM
1. Introduction Superaustenitic stainless steels are widely used in applications where increased pitting corrosion resistance is needed, such as chemical, gas, oil, pulp and paper industries, marine and offshore applications, structure material for heat exchangers, piping or desalination [1–4]. In addition to superior corrosion resistance, these steels combine enhanced mechanical properties, formability and weldability, within a wide temperature range, when compared to conventional austenitic stainless steels or Ni-based superalloys, which were traditionally used in such applications [5–10]. Moreover, these grades have a face centered cubic (fcc) microstructure, exhibiting superior impact toughness and can maintain their impact properties down to very low subzero temperatures, enabling them to act as the best candidate material for cryogenic applications [11]. It is very well known that when these heavily alloyed steels are exposed at elevated temperatures (550–1050 1C), they exhibit the formation of several secondary phases, consequently affecting both mechanical properties and corrosion resistance [9,12–18]. The most important intermetallic phase reported to form during
n
Corresponding author. Tel.: þ30 6947812125; fax: þ30 2107722119. E-mail addresses: [email protected] (T. Koutsoukis), [email protected] (A. Redjaı¨mia), [email protected] (G. Fourlaris). 0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.10.066
heat treatment is s-phase. Its formation is caused by decomposition of austenite, when large quantities of Cr and sometimes Mo diffuse within the austenitic matrix and form Cr-rich compounds, which are responsible for the deterioration in the overall corrosion resistance of the steel. When formed in large volume fraction, the hard and brittle s-phase turns the ductile austenitic steel to more brittle. Other very common secondary phases formed during the ageing in stainless steels could be carbides, nitrides, w-phase or Laves phase. A recent work made by Liu [19], summarizes all the possible phase transformations that can occur in all types of steels. However, there are many questions concerning the formation conditions and mechanism of these secondary phases, especially in superaustenitic stainless steel grades. Although there are enough data for s-phase, yet there are very limited information, for instance, for the Laves phase, which was found in a previous study to form in the microstructure of such steels [20]. The purpose of this work is to study the various phase transformations that take place following ageing and to correlate with the mechanical behavior of superaustenitic stainless steels.
2. Materials and methods The steels used in the present study were hot rolled grades, namely S32654 and S31254 (UNS designation), with 2.27 mm and
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Table 1 Elemental composition (wt%) of S31254 and S32654. Elemental composition
S31254 S32654
C
Si
Mn
P
S
Cr
Ni
Mo
N
Cu
Ti
0.012 0.013
0.36 0.24
0.47 3.43
0.019 0.021
0.001 0.001
20.02 24.19
18.16 21.58
5.98 7.24
0.214 0.497
0.65 0.38
0.001 0.001
3.20 mm in thickness, respectively. S32654 is one of the most heavily alloyed superaustenitic stainless steel, whilst S31254 one of the most frequently used superaustenitic grade. The composition (wt%) of the steel grades is shown in Table 1. S32654 is higher alloyed in Cr, Ni, Mo, N and Cu than S31254. The specimens were mechanically cut to pieces of 10 mm by 20 mm and subjected to isothermal heat treatments within the temperature range of 650–950 1C, with a step of 100 1C, in atmospheric conditions (air), for ageing times varying between 0.5 h and 3000 h, followed by water quenching at room temperature. Exposure to atmospheric conditions was chosen in order to simulate the conditions in service of these grades. As a consequence, oxidation of specimens occurred during 240 h or 1000 h ageing at the temperatures of 950 1C or 850 1C, respectively. Metallographic preparation for scanning electron microscopy (SEM) examination of the specimens, employing a Philips XL30 at 20 kV, was performed involving standardized metallographic mounting and polishing techniques. Energy Dispersive X-ray Microanalysis (EDS) was used in order to analyze the elemental composition of the various secondary phases. Etching of the specimens was performed with 10% oxalic acid in deionized water at 5 V, for etching times varying between 0.5 s and 20 s. Thin foils for transmission electron microscopy (TEM) were prepared following conventional TEM specimen preparation methods. The specimens were mechanically ground down to the lowest possible thickness (20–10 mm) and then electro polished at 40 V in a solution of 5% perchloric acid in 95% II-butoxyethanol, employing a Struers Tenupol twin-jet unit. Specimen examination was performed by a Philips CM12 microscope operated at 120 kV. Diffraction patterns were obtained in the selected-area electron diffraction (SAED) mode with a parallel incident beam and also in the microdiffraction mode, using convergent beam electron diffraction (CBED) with a nearly parallel electron beam focused on a very small area of the thin foil (10–0 nm). Following examination via electron microscopy the results were summarized and the time–temperature–precipitation (TTP) curves were drawn. Vickers hardness testing (HV10, 10 kg loading force) and preliminary tensile testing, according to the ASTM E8M-04 standard, were performed in order to correlate the evolution of the microstructural changes with mechanical properties of the steels. At least 12 measurements per specimen were taken for Vickers hardness testing while three specimens were examined under the same conditions for tensile testing and standard deviations were calculated. Following tensile testing the type of the fracture was determined via SEM examination. The modulus of toughness (UT) was also calculated in order to evaluate the effect of ageing on the variation of toughness. In the present study, the condition where steels were not subjected to any kind of heat treatment is termed as ‘‘As Reference’’ (AR) and is equivalent to the ‘‘As Received’’ condition. All the numerical results were statistically analyzed with the use of the two-way analysis of variance (ANOVA) method, at a ¼0.05 level of significance, employing the SigmaPlot software. Statistically significant differences between groups were found and were determined by the use of Holm–Sidak’s procedure.
3. Experimental results 3.1. Secondary phase precipitation Following isothermal ageing, all specimens were examined in the SEM for a preliminary evaluation of the microstructure. It was possible to distinguish between two different types of precipitates, while imaging, using backscattered mode, based on the difference in atomic number of the elements present in the two species of precipitates when compared to the austenitic matrix. EDS microanalysis showed that some were richer in Cr and the others richer in Mo. Precipitation takes place very soon (less than 1 h) at 950 1C, while much longer periods are required when decreasing the temperature down to 650 1C. As shown in a previous study [20], following careful TEM observation, at least four different types of precipitates were observed. These four identified secondary phases are s-phase, Laves phase, w-phase and the b-Cr2N nitride, with the volume fraction of the two former being significantly greater than the volume fraction of either w-phase or b-Cr2N nitride, in both steel grades. There are different modes of electron diffraction that can be used for their crystallographic characterization [21], namely Convergent Beam Electron Diffraction (CBED) [22], Large Angle Convergent Beam Electron Diffraction (LACBED) [23], Large Angle Convergent DIFfraction (LACDIF) [24], precession [25] or microdiffraction [26,27]. The microdiffraction, a routine technique, is better suited for the characterization of small precipitates resulting from the decomposition of the matrix. In this study, the microdiffraction procedure that was proposed by Morniroli, Steeds and Redjaı¨mia [26,27] was mobilized. All the notations and the conventions given in reference [26] will be used in this paper also. The crystallographic identification procedure of the decomposition products of the austenitic matrix will be focused only on the two intermetallic phases, namely, s-phase and w-phase.
3.1.1. Intermetallic s-phase precipitation and kinetics The main axes necessary for the crystallographic characterization of s-phase are [001], [100] and [110]. The symmetries recorded along these axes are {(4mm), (4mm)} and {4mm, 4mm} for [001] axis (Fig. 1A) and for the other two axes, [100] and [110], are {(2mm), (2mm)} and {2mm, 2mm}, (Fig. 1B and C). These features indicate that the crystal system of s-phase is tetragonal [26]. The absence of Zero Order Laue Zone (ZOLZ) or First Order Laue Zone (FOLZ) periodicity difference along [001] and [110] axes indicates that there is no glide plane perpendicular to these two directions and the individual extinction symbol could be either ‘‘P-..’’ or ‘‘P..-’’. On the other hand, the periodicity difference between the ZOLZ and FOLZ nets reveals the presence of diagonal ‘‘n’’ glide plane perpendicular to [100], leading to the individual partial extinction symbol ‘‘P.n.’’. The addition of the three individual partial extinction symbols indicates that the partial extinction symbol is ‘‘P-n-’’. The latter is in agreement with three space groups, each of them belongs to a point group. By considering the partial extinction symbol and the {(2mm), 2mm} ideal symmetries of the microdiffraction patterns recorded along [100] and
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Fig. 1. Diffraction patterns recorded along (A) [001], (B) [100] and (C) [110] for s-phase.
[110] (Fig. 1), it can be concluded, without any ambiguity, that the intermetallic s-phase is crystallized in the tetragonal system and belongs to the space group ‘‘P 42/m n m’’ (space group no.136) or ‘‘P 42/m 21/n 2/m’’, in its full notation. In order to identify s-phase, clear electron diffraction patterns were able to be taken as s-phase particles exhibit either few or no defects (Fig. 2A). No specific orientation relationships of s-phase and the austenitic matrix were identified, in agreement to previous studies [1,19]. The nucleation sites of s-phase are growing both intergranularly and intragranularly [28], with particles located intergranularly having a faster coarsening rate, reaching up to longer sizes (up to around 10 mm in diameter) than those forming intragranularly. The morphology of s-phase particles was either globular or plate-like. Following EDS microanalysis in the coarser s-phase particles, it was found that its chemical composition was richer in Cr and Mo contents to that of the matrix, especially Cr, and poorer in Ni content, as shown in Table 2. EDS data were used supplementary in the present study, in order to prove Cr partitioning into s-phase particles.
3.1.2. Intermetallic w-phase precipitation and kinetics The main axes necessary for the crystallographic characterization of w-phase are [001], [111] and [011]. The symmetries recorded along the [001] zone axis are {(4mm), (4mm)} and {4mm, 2mm} (Fig. 3A) and those recorded along [1¯11] zone axis are {(6mm), (6mm)} and {3m, 3m} (Fig. 3B). For the zone axis [110], the recorded symmetries are {(2mm), (2mm)} and {2mm, 2mm} (Fig. 3C). These features indicate that the crystal system of the w-phase is cubic [26]. The absence of ZOLZ/FOLZ periodicity difference indicates that there is no glide plane perpendicular to the [001] direction and that the partial extinction symbol is ‘‘I-..’’ or ‘‘F-..’’. The absence of ZOLZ/FOLZ periodicity difference indicates that there is no glide plane perpendicular to the [011] direction and that the partial extinction symbol is ‘‘I..-’’. By combining the information, it is quite clear that the absence of ZOLZ/FOLZ periodicity along these two zone axes indicates that w-phase is crystallized in a cubic system with a partial extinction symbol ‘‘I—’’. The latter is in agreement with seven space groups, each of them belongs to a point group. By considering the partial extinction symbol and the {(4mm), 2mm} ideal symmetries of the microdiffraction patterns recorded along [001] direction, it can be concluded, without any ambiguity, that the intermetallic w-phase is crystallized in the cubic system and belongs to the space group ‘‘I4¯3m’’ (space group no. 217). The volume fraction of w-phase was found to be very small at any heat treated condition. Its size and morphology are very similar to that of s-phase, which made their distinction challenging. Only a few particles were found randomly distributed intragranularly, with sizes ranging to a few microns. w-phase was observed at the specimens heat treated at the higher temperatures of this study (850 1C or
950 1C) and for ageing times over 24 h in both steels. Nevertheless, the limited volume fraction of crystal defects observed in w-phase particles allowed the successful determination of the required zone axis. Two kinds of orientation relationships (OR) were derived between w-phase and the austenitic matrix (g). One was found to be the Nishiyama–Wassermann, where: (110)w:(111)g,[1¯10]w: [1¯21¯]g, and [001]w: [1¯01]g, which have been also reported in previous studies [1,19,29]. The second one is at 31 close to the cubic-on-cube orientation relationship, along zone axis [011]w: [011]g and [1¯11]w: [1¯11]g, where: (1¯11)w31:(1¯11)g, [110]w31:[110]g and [011¯]w31:[011¯]g [20]. Fig. 2C shows a typical precipitate of w-phase.
3.1.3. Intermetallic Laves phase precipitation and kinetics Laves phase particles were observed in large volume fraction, similar to the volume fraction of s-phase, but were found to form prior to s-phase precipitates, during ageing at the lower temperatures of this study (650 1C or 750 1C). Laves phase has a stoichiometry of A2B, is primarily crystallized in the hexagonal system and belongs to the space group ‘‘P 63/m m c’’ (space group no. 194), or ‘‘P 63/m 2/m 2/c’’, in its full notation. It is difficult to identify Laves phase and the major reason is the dense network of defects present in particles, as shown in Fig. 2D. When the electron beam is concentrated on a particle of Laves phase, in addition to the desired reflected crystal planes, diffusion intensity lines (DIL) are also formed due to plane defects, resulting in a complex diffraction pattern that is very difficult to index. The dense network of defects is often observed in Laves phase formed in stainless steels, rather than other alloy category. What is more, Laves phase could either form in the hexagonal (C14, C36) or sometimes in the cubic (C15) system, having a transition temperature that ranges in a different way among various alloy systems [30]. In the literature, there are very few, if any, clear diffraction patterns of the Laves phase observed in stainless steels [19,31]. It is necessary to find an area free of defects on a Laves phase particle in order to receive such an electron diffraction pattern and this was accomplished and pointed out in Fig. 4. Analysis of the different recorded electron diffraction patterns indicates that the Laves phase develops, with the austenitic matrix, the following orientation relationship: ([0001]L:[001]g), (200)g:(2¯110)L and (020)g:(011¯0)L. The nucleation sites of Laves phase particles were primarily observed intragranularly and its morphology varied, being either polyhedral-like or plate-like, sometimes needle-like and in much fewer cases globular. For the last morphology, the diameter was measured between a few tenths of a nanometer up to about 2 mm. Laves phase precipitates were always observed uniformly dispersed within the whole specimen. Laves phase is the first secondary phase to form, following ageing at all temperatures, followed by the formation of s-phase. What is also interesting is that almost no particles of Laves phase were observed following ageing at 950 1C for 240 h, implying full transformation of this phase
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Fig. 2. Bright field TEM micrograph of S31254 showing (A) Intragranular s-phase particle (2000 h, 750 1C). Note the twins induced by the hard s-phase. (B) Alternation between Laves phase and b-Cr2N precipitates along grain boundary (240 h, 750 1C). (C) w-Phase particle (48 h, 950 1C) and (D) Laves phase particle (24 h, 850 1C). Note the dense network of defects.
Table 2 Elemental composition range (wt%) of s-phase, Laves phase and the austenitic matrix measured via EDS microanalysis. Each range includes 10 measurements per steel from both S31254 and S32654. Element
Phase a
Cr Mo Ni
s
24.8–35.3 13.5–22.5 6.8–19.5
a
s-phase.
b
Austenitic matrix.
Laves
b
19.2–20.5 21.5–27.4 10.6–12.3
20.7–22.5 4.7–6.9 24.7–25.2
c (around precipitates)
b
c (AR)
20–24.2 6–7.2 18.2–21.6
under these conditions. Laves phase is rich in Mo, as measured via EDS (Table 2). EDS data were used in addition to the TEM results, in order to prove the partitioning of Mo into Laves phase particles.
3.1.4. b-Cr2N precipitation and kinetics A fourth species of precipitation was also observed to form both intragranularly or intergranularly and based on the electron diffraction patterns it is believed to be the Cr2N nitride, also known as
b-Cr2N nitride. The b-Cr2N nitride has a trigonal crystal structure ‘‘P3¯1m’’ (space group no. 162), described with a hexagonal subcell, while there are various orientation relationships with the austenitic matrix, with the most well known being the Presser–Silcock, where [0001]b:[110]g and [21¯1¯0]b:[11¯2]g [19]. Nitrides are prone to form in these steels following heat treatment because of the high amounts of N in the alloy composition [2,15,16,31–33]. What is interesting about the nitride formation behavior is that it was always found to nucleate adjacent to Laves phase particles in the austenitic matrix, as shown in Fig. 2B. Morphologically, the b-Cr2N nitride has an unspecified shape, differing to the more longitudinal shaped nitrides reported to form in the literature in similar steels [33,34], while its maximum size reaches up to about 400 nm in diameter. The g-matrix lattice parameter, ag ¼0.359 nm, characterized by X-ray diffraction, is used to calibrate the electron diffraction patterns to deduce the lattice parameters of the four secondary phases. The results are as follows: a¼0.8799 nm, c¼0.4544 nm, c/a¼0.516, for s-phase, a¼0.894 nm for w-phase, a¼0.4744 nm, c¼0.7725 nm, for Laves phase and a¼0.4795 nm, c¼0.4469 nm, for b-Cr2N, and the crystallographic data are in agreement with previous studies [1,19]. Summarizing all the above mentioned information about the nucleation and coarsening of the various precipitates observed within the experimental conditions of this study, the TTP diagrams were
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Fig. 3. Diffraction patterns recorded along (A) [001], (B) [1¯11] and (C) [011] for s-phase. The mirrors (m) crossed in FOLZ on Fig. 3A indicate that these mirrors are destroyed and the symmetry is lost in the FOLZ. The WP 4 mm net symmetry is then reduced to the 2 mm ideal symmetry.
drawn, as shown in Fig. 5, which were preliminary presented by the authors in a previous study [20]. Furthermore, although, Heino[31] and Heino et al. [32], and Lee et al. [14,15,33] have proven the formation of various carbides or other secondary phases, like R-phase, in these steels, none of these were identified in the present study. 3.2. Mechanical behavior The Vickers hardness curves of S31254 and S32654 steels were drawn, as shown in Fig. 6. Comparing the curves in Fig. 6A with those in Fig. 6B, it is shown that both steel grades exhibit similar hardness evolution at each temperature, with the hardness curves of S32654 shifted to higher values than those of S31254. Coarsening of precipitates over a critical size and dispersion occurs at shorter ageing times at 950 1C, where precipitation kinetics is faster. On the contrary, at 650 1C the hardness values seem to keep increasing even following 3000 h of ageing, implying that the peak on the curve might have not been reached as yet. Yield stress (s0.2%) and ultimate tensile stress (UTS) values of the AR specimens are lower than the nominal yield stress and UTS values, respectively. As shown in Fig. 7A, the yield stress and the UTS have either maintained (UTS values at 650 1C) or increased their values in all experimental conditions, compared to the AR values. The highest value was measured at 750 1C following 240 h of ageing. In addition, the elongation (e%) values decrease to the ageing time, while the AR elongation value is significantly higher than the nominal value (Fig. 7B). As an exception, no statistical difference was found between the elongation values measured at 950 1C, for either 120 h or 240 h. Fractography via SEM, revealed a combination of both ductile and brittle fracture in all specimens studied, with a more brittle behavior taking place following ageing at higher ageing temperatures and times. The modulus of toughness (UT) was calculated as the total area under the stress–strain curve [35]. Calculating the changes in UT, caused by precipitation, it is possible to evaluate the overall behavior of toughness versus ageing temperature and time. At any temperature, comparing the values obtained between 24 h and 240 h of ageing, the toughness decreases (Fig. 7C). However, no statistically significant differences were found between the 120 h of ageing and 24 h or 240 h.
4. Discussion 4.1. Precipitation of secondary phases in austenite The ageing temperature and time of the formation of the precipitates seem to be strictly dependent on the elemental composition of the steels. Even slight changes in the elemental composition of these steels are often the reason for somewhat different results between the various studies, on the same steel grades [14]. Small changes in the composition in Mo ( 71%) [31],
and even smaller in N ( 70.1%) [2,10,36] are sufficient to promote or retard the formation of different secondary phases in the same ageing times and temperatures, or similar secondary phases in different ageing times and temperatures. As shown in Fig. 5, precipitation of s-phase occurs at shorter ageing times in the most heavily alloyed with Cr and Mo S32654, when compared to S31254. In terms of ageing time, the first intermetallic phase that forms in both the steel grades, is the Laves phase, during ageing at either 650 1C or 750 1C, whilst at 950 1C Laves phase and s-phase nucleate almost at the same time in both steel grades, which is in agreement to previous studies [16,31]. Nevertheless, as Laves phase is a metastable phase, following ageing at 950 1C for 240 h very few Laves phase precipitates were observed. At the same time, the volume fraction of s-phase was increased, forming coarser particles, which implies either the transformation of Laves phase into s-phase, or the dissolution of Laves phase into the austenitic matrix followed by the formation of s-phase, or even both. Precipitation kinetics is much more sluggish at 850 1C, 750 1C and even more at 650 1C. Once Laves phase is formed, the same phase transformation is expected to occur at these temperatures, too, but at longer ageing times. It is believed that the dissolution of Laves phase is most possible to occur, as the relation between some of the crystallographic directions and planes of the hexagonal crystal system (Laves phase) and the fcc system of austenite is stronger than that of the tetragonal system of s-phase and the fcc system of austenite. At all the heat treatments performed in both steel grades the b-Cr2N nitride was always found adjacent to Laves phase particles and, given the fact that the Laves phase precipitates nucleate prior to the b-Cr2N nitride, the b-Cr2N nitride formation seems to be favored by the presence of Laves phase. Diffusion of Mo into austenite and precipitation of Laves phase causes the depletion of Mo in the surrounding matrix and promotes the formation of Cr-rich phases, such as s-phase and the b-Cr2N, turning the austenitic matrix richer in Ni as shown in Table 2. This could be an explanation for the formation of b-Cr2N particles adjacent to Laves phase particles (Fig. 2B). The b-Cr2N nitrides often form a continuous network of alternating precipitates along grain boundaries, which was also reported in previous studies [31]. Moreover, the initiation of the b-Cr2N nitride formation takes place at shorter ageing times in S32654 than in S31254, because of the higher content in N (Fig. 5). The cellular type of b-Cr2N was not observed in this study [1,33]. The b-Cr2N nitrides are more plate-like when observed intragranularly or adjacent to other types of precipitates (Fig. 2B) and more longitudinal when found single in the austenitic matrix and this is due to the faster diffusion of Cr in the direction of [0001]b:[111]g [19]. In the present study only a few longitudinal b-Cr2N precipitates were found alone in the austenitic matrix with such a shape. A few precipitates of w-phase were also found in specimens aged at 850 1C or 950 1C, a temperature range which is not very common for this phase transformation in stainless steels [29,37].
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Fig. 5. TTP diagrams for (A) S31254 and (B) S32654 [20].
1000 h of ageing at 850 1C, some w-phase precipitates were detected around the predominant s-phase precipitates. 4.2. Correlation of precipitation with mechanical behavior
Fig. 4. (A) Bright field TEM micrograph of a Laves phase precipitate observed following ageing of 1000 h at 850 1C in S31254 steel and (B) electron diffraction pattern recorded along [0001]L//[001]g from the particle of Fig. 4A, where (200)g: (2¯110)L and (020)g: [011¯0]L.
This phase has similar morphology to s-phase and it is often hard to distinguish them without the use of electron diffraction techniques. In addition, minimum or no defects were observed in the precipitates of w-phase, which is unusual to other types of stainless steels, such as duplex stainless steels [29], where w-phase is usually observed at lower temperatures with a lot of defects in its crystal structure. The volume fraction of w-phase does not increase as much as that of s-phase, but still, after
Within the initial ageing hours at either 850 1C or 950 1C the size of the precipitates remains at submicron values and precipitates form in a uniform dispersion, which results in an increase of hardness values. At 850 1C, hardness values increase up to a maximum following ageing for 120 h (Fig. 6), in both steel grades, where the combination of size and dispersion of Laves phase, s-phase and also b-Cr2N nitrides is optimized favoring hardness. In addition to precipitation hardening, another factor that assists the increase of hardness is hardening via second phase formation, since the volume fraction of the hard s-phase has been significantly increased and the size of the precipitates reach up to 10 mm. Following 120 h of ageing, partial dissolution or transformation of Laves phase to s-phase and coarsening of s-phase precipitates results to an overall decrease in hardness. The same behavior was observed during ageing at 950 1C, but coarsening of both Laves phase and s-phase and dissolution of Laves phase occurs at shorter time intervals. Moreover, increased amounts of Mo and Cr in S32654 lead to the formation of higher volume fraction of s-phase and Laves phase, than those formed in S31254. The decrease in hardness at 950 1C between 96 h and 120 h is statistically significant (Po0.001) and could be attributed to the change in hardness mechanism. It is the time period where the volume fraction of Laves phase decreases intensely and Mo dissolves to the austenitic matrix. Following that time interval, s-phase is primarily responsible for the further evolution of hardness.
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Fig. 6. Hardness evolution curves (HV) for (A) S31254 and (B) S32654.
As precipitation kinetics is slower and also different at 750 1C, during the first 24 h of ageing no s-phase was formed in S31254 and a few particles of s-phase were formed in S32654, with Laves phase being the predominant phase, but still forming in dispersed precipitates primarily along the grain boundaries. These precipitates may cause a small amount of precipitation hardening, but possible stress relief at 750 1C overcomes precipitation hardening and causes a decrease in overall hardness values. Also, as shown in Fig. 2A, mechanical twins were observed via TEM examination, caused by the growth of s-phase, which indicates that s-phase is hard enough to generate stresses at the s/g interface causing twinning of austenite. On the contrary, precipitation hardening during the initial ageing hours at 850 1C or 950 1C prevailed, resulting in a very fast increase of hardness values. The combination of different mechanisms that result into an overall decrease of hardness in both steel grades is better shown schematically by the ageing curves at 650 1C (Fig. 6). Up to 240 h of ageing hardness values lie below the AR value (P o0.001) while no statistically significant difference was found between the AR value and those measured at 500 h of ageing. This implies that original hardness value is achieved following 500 h of ageing. Precipitation kinetics at 650 1C is very slow with Laves phase being the predominant phase identified. Laves phase precipitates nucleate and coarsen at a very slow rate and hardness values increase over the AR values, only following 500 h of ageing in both steel grades. Few s-phase precipitates and b-Cr2N nitrides were observed at the specimens exposed for 3000 h of ageing, which is the longest ageing time for the present study. This substantially means that only Laves phase is involved in the evolution of mechanical behavior at 650 1C.
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As shown in Fig. 7A both UTS and yield stress values have been increased under all heat treatment conditions when compared to the AR values. The behavior at 650 1C is an exception, where both yield stress and UTS have either maintained or decreased their values, compared to the AR state, because no precipitation of the s-phase was observed and the volume fraction of Laves phase precipitates was minor. The overall strength of the steel was thus maintained or decreased because of possible stress relief that could take place during the initial ageing hours. Tensile behavior is similar to hardness behavior, meaning that when hardness values increase, UTS and yield stress values increase and vice versa. The optimum combination of Laves phase and s-phase dispersion and particle size leads to the maximum values of UTS and yield stress which correspond to the intermediate, for this study, at temperatures of 750 1C or 850 1C. Following heat treatments at 650 1C, slow precipitation kinetics of Laves phase and the absence of other types of secondary phases result in lower UTS and yield stress values, compared to the other ageing temperatures (Fig. 7A). Increased values of UTS and yield stress have turned the steel less ductile, as can be concluded by Fig. 7B. At 750 1C or 850 1C, the coarse s-phase particles and a uniform dispersion of Laves phase particles are responsible for a reduction in ductility values, compared to the AR value. Following ageing at 850 1C for 240 h, the elongation of S31254 is reduced by more than 65% and this is primarily attributed to the large, brittle precipitates of s-phase, but also to the precipitates of Laves phase and to a lesser degree to that of the dispersed b-Cr2N nitrides. The only experimental condition where elongation maintained the AR value is ageing at 650 1C for 24 h, because the small volume fraction of Laves phase precipitates has no major effect on ductility. However, considering the minimum nominal value for ductility of S31254 (Fig. 7B), it could be stated that the steel essentially maintains its elongation requirements following ageing at 650 1C for up to 240 h. Furthermore, the change in the modulus of toughness (UT), which is the amount of the work per unit volume versus ageing time and temperature, was also studied. The UT values are calculated by the total area under the stress–strain curve and should not be confused or compared to the impact of toughness values, which have been measured in previous studies [38,39], because it includes not only the energy absorbed during fracture but also the energy absorbed during plastic deformation. The evolution of UT versus ageing time and temperature (Fig. 7C) has similar behavior to the evolution of elongation (Fig. 7B), with all UT values calculated below the AR value of the steel in all experimental conditions. As an exception, toughness is essentially maintained following ageing at 650 1C for 24 h, as no precipitations were observed and no statistical significance was calculated for this heat treatment (P 4 0.05). Laves phase and s-phase precipitates dramatically affected the UT values, decreasing the values more than 30% of the AR values, following ageing at 850 1C for 240 h. This decrement was reasonable as elongation was found to decrease and so did the total area below the stress–strain curve, even if UTS was found to increase. As an overall comment for the present study it can be derived that Laves phase and/or s-phase are primarily responsible for all the changes in the mechanical properties of the aged superaustenitic stainless steels, while minor effect could be attributed to either w-phase or the b-Cr2N nitride. Following ageing above 750 1C both steel grades suffer significant changes in the mechanical properties, deteriorating their ductility but improving their strength, while it is much safer to expose these steel grades below 650 1C, where precipitation kinetics is much slower and mechanical properties are substantially improved, compared to higher temperatures.
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Fig. 7. Evolution of (A) yield stress s0.2% (MPa) and UTS (MPa), (B) elongation e (%) and (C) modulus of toughness UT (TPa) versus ageing time t (h) for S31254 (nom.: nominal value, AR: as measured in the as received condition).
5. Conclusions
At least four different types of precipitates, s-phase, w-phase,
Laves phase and b-Cr2N nitride, were identified following ageing of S32654 and S31254 within the temperature range of 650– 950 1C. At all temperatures, Laves phase is the first secondary phase forming, followed by s-phase and then at longer ageing times w-phase and the b-Cr2N nitride also form. Secondary phases form at earlier ageing times in the most heavily alloyed S32654. Full transformation of Laves phase to s-phase and/or its dissolution in the austenitic matrix takes place for time intervals over 240 h at an ageing temperature of 950 1C. Precipitation of primarily Laves phase and s-phase resulted into an increase of the hardness values at all temperatures in both the steel grades. Secondary phase precipitation resulted into an increase of both yield stress and UTS while elongation and toughness were decreased in both steel grades.
Appendix A. Supplementary materials Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.msea.2012.10.066.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Phase transformations and mechanical properties in heat treated superaustenitic stainless steels T. Koutsoukis a,n, A. Redjaı¨mia b, G. Fourlaris a a b
Laboratory of Physical Metallurgy, National Technical University of Athens, 9 Heroon Polytechniou Street, 15780 Athens, Greece Universite´ de Lorraine, Institut Jean Lamour, UMR-CNRS 7198, Nancy F-54042, France
a r t i c l e i n f o
abstract
Article history: Received 27 June 2012 Received in revised form 16 October 2012 Accepted 17 October 2012 Available online 26 October 2012
A microstructure–properties relationship study in two superaustenitic stainless steels (S31254 and S32654) was carried out, following exposure at elevated temperatures for various ageing times. Due to high temperature ageing, most stainless steel grades suffer the formation of various precipitates, directly affecting their properties. The full characterization of those precipitates and the correlation with the mechanical behavior of the steels is the primary aim of this study. Samples of the steel grades studied, were exposed to isothermal heat treatments within the temperature range of 650–950 1C, for ageing times varying between 0.5 h and 3000 h, followed by water quenching at room temperature. Microstructural examination indicated the formation of four different secondary phases, sigma phase (s), chi phase (w), Laves phase and b-Cr2N nitride, which were characterized by transmission electron microscopy (TEM) and electron diffraction. The results obtained permitted the construction of the time–temperature–precipitation (TTP) plots. In addition, tensile and Vickers hardness testing were utilized and the modulus of toughness was calculated. The kinetics of the formation of various precipitates with increasing temperature and aging duration was also observed. It was found that various precipitates had a significant effect on all mechanical properties studied. & 2012 Elsevier B.V. All rights reserved.
Keywords: Superaustenitic stainless steel Ageing Precipitation Laves phase Mechanical properties TEM
1. Introduction Superaustenitic stainless steels are widely used in applications where increased pitting corrosion resistance is needed, such as chemical, gas, oil, pulp and paper industries, marine and offshore applications, structure material for heat exchangers, piping or desalination [1–4]. In addition to superior corrosion resistance, these steels combine enhanced mechanical properties, formability and weldability, within a wide temperature range, when compared to conventional austenitic stainless steels or Ni-based superalloys, which were traditionally used in such applications [5–10]. Moreover, these grades have a face centered cubic (fcc) microstructure, exhibiting superior impact toughness and can maintain their impact properties down to very low subzero temperatures, enabling them to act as the best candidate material for cryogenic applications [11]. It is very well known that when these heavily alloyed steels are exposed at elevated temperatures (550–1050 1C), they exhibit the formation of several secondary phases, consequently affecting both mechanical properties and corrosion resistance [9,12–18]. The most important intermetallic phase reported to form during
n
Corresponding author. Tel.: þ30 6947812125; fax: þ30 2107722119. E-mail addresses: [email protected] (T. Koutsoukis), [email protected] (A. Redjaı¨mia), [email protected] (G. Fourlaris). 0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.10.066
heat treatment is s-phase. Its formation is caused by decomposition of austenite, when large quantities of Cr and sometimes Mo diffuse within the austenitic matrix and form Cr-rich compounds, which are responsible for the deterioration in the overall corrosion resistance of the steel. When formed in large volume fraction, the hard and brittle s-phase turns the ductile austenitic steel to more brittle. Other very common secondary phases formed during the ageing in stainless steels could be carbides, nitrides, w-phase or Laves phase. A recent work made by Liu [19], summarizes all the possible phase transformations that can occur in all types of steels. However, there are many questions concerning the formation conditions and mechanism of these secondary phases, especially in superaustenitic stainless steel grades. Although there are enough data for s-phase, yet there are very limited information, for instance, for the Laves phase, which was found in a previous study to form in the microstructure of such steels [20]. The purpose of this work is to study the various phase transformations that take place following ageing and to correlate with the mechanical behavior of superaustenitic stainless steels.
2. Materials and methods The steels used in the present study were hot rolled grades, namely S32654 and S31254 (UNS designation), with 2.27 mm and
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Table 1 Elemental composition (wt%) of S31254 and S32654. Elemental composition
S31254 S32654
C
Si
Mn
P
S
Cr
Ni
Mo
N
Cu
Ti
0.012 0.013
0.36 0.24
0.47 3.43
0.019 0.021
0.001 0.001
20.02 24.19
18.16 21.58
5.98 7.24
0.214 0.497
0.65 0.38
0.001 0.001
3.20 mm in thickness, respectively. S32654 is one of the most heavily alloyed superaustenitic stainless steel, whilst S31254 one of the most frequently used superaustenitic grade. The composition (wt%) of the steel grades is shown in Table 1. S32654 is higher alloyed in Cr, Ni, Mo, N and Cu than S31254. The specimens were mechanically cut to pieces of 10 mm by 20 mm and subjected to isothermal heat treatments within the temperature range of 650–950 1C, with a step of 100 1C, in atmospheric conditions (air), for ageing times varying between 0.5 h and 3000 h, followed by water quenching at room temperature. Exposure to atmospheric conditions was chosen in order to simulate the conditions in service of these grades. As a consequence, oxidation of specimens occurred during 240 h or 1000 h ageing at the temperatures of 950 1C or 850 1C, respectively. Metallographic preparation for scanning electron microscopy (SEM) examination of the specimens, employing a Philips XL30 at 20 kV, was performed involving standardized metallographic mounting and polishing techniques. Energy Dispersive X-ray Microanalysis (EDS) was used in order to analyze the elemental composition of the various secondary phases. Etching of the specimens was performed with 10% oxalic acid in deionized water at 5 V, for etching times varying between 0.5 s and 20 s. Thin foils for transmission electron microscopy (TEM) were prepared following conventional TEM specimen preparation methods. The specimens were mechanically ground down to the lowest possible thickness (20–10 mm) and then electro polished at 40 V in a solution of 5% perchloric acid in 95% II-butoxyethanol, employing a Struers Tenupol twin-jet unit. Specimen examination was performed by a Philips CM12 microscope operated at 120 kV. Diffraction patterns were obtained in the selected-area electron diffraction (SAED) mode with a parallel incident beam and also in the microdiffraction mode, using convergent beam electron diffraction (CBED) with a nearly parallel electron beam focused on a very small area of the thin foil (10–0 nm). Following examination via electron microscopy the results were summarized and the time–temperature–precipitation (TTP) curves were drawn. Vickers hardness testing (HV10, 10 kg loading force) and preliminary tensile testing, according to the ASTM E8M-04 standard, were performed in order to correlate the evolution of the microstructural changes with mechanical properties of the steels. At least 12 measurements per specimen were taken for Vickers hardness testing while three specimens were examined under the same conditions for tensile testing and standard deviations were calculated. Following tensile testing the type of the fracture was determined via SEM examination. The modulus of toughness (UT) was also calculated in order to evaluate the effect of ageing on the variation of toughness. In the present study, the condition where steels were not subjected to any kind of heat treatment is termed as ‘‘As Reference’’ (AR) and is equivalent to the ‘‘As Received’’ condition. All the numerical results were statistically analyzed with the use of the two-way analysis of variance (ANOVA) method, at a ¼0.05 level of significance, employing the SigmaPlot software. Statistically significant differences between groups were found and were determined by the use of Holm–Sidak’s procedure.
3. Experimental results 3.1. Secondary phase precipitation Following isothermal ageing, all specimens were examined in the SEM for a preliminary evaluation of the microstructure. It was possible to distinguish between two different types of precipitates, while imaging, using backscattered mode, based on the difference in atomic number of the elements present in the two species of precipitates when compared to the austenitic matrix. EDS microanalysis showed that some were richer in Cr and the others richer in Mo. Precipitation takes place very soon (less than 1 h) at 950 1C, while much longer periods are required when decreasing the temperature down to 650 1C. As shown in a previous study [20], following careful TEM observation, at least four different types of precipitates were observed. These four identified secondary phases are s-phase, Laves phase, w-phase and the b-Cr2N nitride, with the volume fraction of the two former being significantly greater than the volume fraction of either w-phase or b-Cr2N nitride, in both steel grades. There are different modes of electron diffraction that can be used for their crystallographic characterization [21], namely Convergent Beam Electron Diffraction (CBED) [22], Large Angle Convergent Beam Electron Diffraction (LACBED) [23], Large Angle Convergent DIFfraction (LACDIF) [24], precession [25] or microdiffraction [26,27]. The microdiffraction, a routine technique, is better suited for the characterization of small precipitates resulting from the decomposition of the matrix. In this study, the microdiffraction procedure that was proposed by Morniroli, Steeds and Redjaı¨mia [26,27] was mobilized. All the notations and the conventions given in reference [26] will be used in this paper also. The crystallographic identification procedure of the decomposition products of the austenitic matrix will be focused only on the two intermetallic phases, namely, s-phase and w-phase.
3.1.1. Intermetallic s-phase precipitation and kinetics The main axes necessary for the crystallographic characterization of s-phase are [001], [100] and [110]. The symmetries recorded along these axes are {(4mm), (4mm)} and {4mm, 4mm} for [001] axis (Fig. 1A) and for the other two axes, [100] and [110], are {(2mm), (2mm)} and {2mm, 2mm}, (Fig. 1B and C). These features indicate that the crystal system of s-phase is tetragonal [26]. The absence of Zero Order Laue Zone (ZOLZ) or First Order Laue Zone (FOLZ) periodicity difference along [001] and [110] axes indicates that there is no glide plane perpendicular to these two directions and the individual extinction symbol could be either ‘‘P-..’’ or ‘‘P..-’’. On the other hand, the periodicity difference between the ZOLZ and FOLZ nets reveals the presence of diagonal ‘‘n’’ glide plane perpendicular to [100], leading to the individual partial extinction symbol ‘‘P.n.’’. The addition of the three individual partial extinction symbols indicates that the partial extinction symbol is ‘‘P-n-’’. The latter is in agreement with three space groups, each of them belongs to a point group. By considering the partial extinction symbol and the {(2mm), 2mm} ideal symmetries of the microdiffraction patterns recorded along [100] and
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Fig. 1. Diffraction patterns recorded along (A) [001], (B) [100] and (C) [110] for s-phase.
[110] (Fig. 1), it can be concluded, without any ambiguity, that the intermetallic s-phase is crystallized in the tetragonal system and belongs to the space group ‘‘P 42/m n m’’ (space group no.136) or ‘‘P 42/m 21/n 2/m’’, in its full notation. In order to identify s-phase, clear electron diffraction patterns were able to be taken as s-phase particles exhibit either few or no defects (Fig. 2A). No specific orientation relationships of s-phase and the austenitic matrix were identified, in agreement to previous studies [1,19]. The nucleation sites of s-phase are growing both intergranularly and intragranularly [28], with particles located intergranularly having a faster coarsening rate, reaching up to longer sizes (up to around 10 mm in diameter) than those forming intragranularly. The morphology of s-phase particles was either globular or plate-like. Following EDS microanalysis in the coarser s-phase particles, it was found that its chemical composition was richer in Cr and Mo contents to that of the matrix, especially Cr, and poorer in Ni content, as shown in Table 2. EDS data were used supplementary in the present study, in order to prove Cr partitioning into s-phase particles.
3.1.2. Intermetallic w-phase precipitation and kinetics The main axes necessary for the crystallographic characterization of w-phase are [001], [111] and [011]. The symmetries recorded along the [001] zone axis are {(4mm), (4mm)} and {4mm, 2mm} (Fig. 3A) and those recorded along [1¯11] zone axis are {(6mm), (6mm)} and {3m, 3m} (Fig. 3B). For the zone axis [110], the recorded symmetries are {(2mm), (2mm)} and {2mm, 2mm} (Fig. 3C). These features indicate that the crystal system of the w-phase is cubic [26]. The absence of ZOLZ/FOLZ periodicity difference indicates that there is no glide plane perpendicular to the [001] direction and that the partial extinction symbol is ‘‘I-..’’ or ‘‘F-..’’. The absence of ZOLZ/FOLZ periodicity difference indicates that there is no glide plane perpendicular to the [011] direction and that the partial extinction symbol is ‘‘I..-’’. By combining the information, it is quite clear that the absence of ZOLZ/FOLZ periodicity along these two zone axes indicates that w-phase is crystallized in a cubic system with a partial extinction symbol ‘‘I—’’. The latter is in agreement with seven space groups, each of them belongs to a point group. By considering the partial extinction symbol and the {(4mm), 2mm} ideal symmetries of the microdiffraction patterns recorded along [001] direction, it can be concluded, without any ambiguity, that the intermetallic w-phase is crystallized in the cubic system and belongs to the space group ‘‘I4¯3m’’ (space group no. 217). The volume fraction of w-phase was found to be very small at any heat treated condition. Its size and morphology are very similar to that of s-phase, which made their distinction challenging. Only a few particles were found randomly distributed intragranularly, with sizes ranging to a few microns. w-phase was observed at the specimens heat treated at the higher temperatures of this study (850 1C or
950 1C) and for ageing times over 24 h in both steels. Nevertheless, the limited volume fraction of crystal defects observed in w-phase particles allowed the successful determination of the required zone axis. Two kinds of orientation relationships (OR) were derived between w-phase and the austenitic matrix (g). One was found to be the Nishiyama–Wassermann, where: (110)w:(111)g,[1¯10]w: [1¯21¯]g, and [001]w: [1¯01]g, which have been also reported in previous studies [1,19,29]. The second one is at 31 close to the cubic-on-cube orientation relationship, along zone axis [011]w: [011]g and [1¯11]w: [1¯11]g, where: (1¯11)w31:(1¯11)g, [110]w31:[110]g and [011¯]w31:[011¯]g [20]. Fig. 2C shows a typical precipitate of w-phase.
3.1.3. Intermetallic Laves phase precipitation and kinetics Laves phase particles were observed in large volume fraction, similar to the volume fraction of s-phase, but were found to form prior to s-phase precipitates, during ageing at the lower temperatures of this study (650 1C or 750 1C). Laves phase has a stoichiometry of A2B, is primarily crystallized in the hexagonal system and belongs to the space group ‘‘P 63/m m c’’ (space group no. 194), or ‘‘P 63/m 2/m 2/c’’, in its full notation. It is difficult to identify Laves phase and the major reason is the dense network of defects present in particles, as shown in Fig. 2D. When the electron beam is concentrated on a particle of Laves phase, in addition to the desired reflected crystal planes, diffusion intensity lines (DIL) are also formed due to plane defects, resulting in a complex diffraction pattern that is very difficult to index. The dense network of defects is often observed in Laves phase formed in stainless steels, rather than other alloy category. What is more, Laves phase could either form in the hexagonal (C14, C36) or sometimes in the cubic (C15) system, having a transition temperature that ranges in a different way among various alloy systems [30]. In the literature, there are very few, if any, clear diffraction patterns of the Laves phase observed in stainless steels [19,31]. It is necessary to find an area free of defects on a Laves phase particle in order to receive such an electron diffraction pattern and this was accomplished and pointed out in Fig. 4. Analysis of the different recorded electron diffraction patterns indicates that the Laves phase develops, with the austenitic matrix, the following orientation relationship: ([0001]L:[001]g), (200)g:(2¯110)L and (020)g:(011¯0)L. The nucleation sites of Laves phase particles were primarily observed intragranularly and its morphology varied, being either polyhedral-like or plate-like, sometimes needle-like and in much fewer cases globular. For the last morphology, the diameter was measured between a few tenths of a nanometer up to about 2 mm. Laves phase precipitates were always observed uniformly dispersed within the whole specimen. Laves phase is the first secondary phase to form, following ageing at all temperatures, followed by the formation of s-phase. What is also interesting is that almost no particles of Laves phase were observed following ageing at 950 1C for 240 h, implying full transformation of this phase
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Fig. 2. Bright field TEM micrograph of S31254 showing (A) Intragranular s-phase particle (2000 h, 750 1C). Note the twins induced by the hard s-phase. (B) Alternation between Laves phase and b-Cr2N precipitates along grain boundary (240 h, 750 1C). (C) w-Phase particle (48 h, 950 1C) and (D) Laves phase particle (24 h, 850 1C). Note the dense network of defects.
Table 2 Elemental composition range (wt%) of s-phase, Laves phase and the austenitic matrix measured via EDS microanalysis. Each range includes 10 measurements per steel from both S31254 and S32654. Element
Phase a
Cr Mo Ni
s
24.8–35.3 13.5–22.5 6.8–19.5
a
s-phase.
b
Austenitic matrix.
Laves
b
19.2–20.5 21.5–27.4 10.6–12.3
20.7–22.5 4.7–6.9 24.7–25.2
c (around precipitates)
b
c (AR)
20–24.2 6–7.2 18.2–21.6
under these conditions. Laves phase is rich in Mo, as measured via EDS (Table 2). EDS data were used in addition to the TEM results, in order to prove the partitioning of Mo into Laves phase particles.
3.1.4. b-Cr2N precipitation and kinetics A fourth species of precipitation was also observed to form both intragranularly or intergranularly and based on the electron diffraction patterns it is believed to be the Cr2N nitride, also known as
b-Cr2N nitride. The b-Cr2N nitride has a trigonal crystal structure ‘‘P3¯1m’’ (space group no. 162), described with a hexagonal subcell, while there are various orientation relationships with the austenitic matrix, with the most well known being the Presser–Silcock, where [0001]b:[110]g and [21¯1¯0]b:[11¯2]g [19]. Nitrides are prone to form in these steels following heat treatment because of the high amounts of N in the alloy composition [2,15,16,31–33]. What is interesting about the nitride formation behavior is that it was always found to nucleate adjacent to Laves phase particles in the austenitic matrix, as shown in Fig. 2B. Morphologically, the b-Cr2N nitride has an unspecified shape, differing to the more longitudinal shaped nitrides reported to form in the literature in similar steels [33,34], while its maximum size reaches up to about 400 nm in diameter. The g-matrix lattice parameter, ag ¼0.359 nm, characterized by X-ray diffraction, is used to calibrate the electron diffraction patterns to deduce the lattice parameters of the four secondary phases. The results are as follows: a¼0.8799 nm, c¼0.4544 nm, c/a¼0.516, for s-phase, a¼0.894 nm for w-phase, a¼0.4744 nm, c¼0.7725 nm, for Laves phase and a¼0.4795 nm, c¼0.4469 nm, for b-Cr2N, and the crystallographic data are in agreement with previous studies [1,19]. Summarizing all the above mentioned information about the nucleation and coarsening of the various precipitates observed within the experimental conditions of this study, the TTP diagrams were
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Fig. 3. Diffraction patterns recorded along (A) [001], (B) [1¯11] and (C) [011] for s-phase. The mirrors (m) crossed in FOLZ on Fig. 3A indicate that these mirrors are destroyed and the symmetry is lost in the FOLZ. The WP 4 mm net symmetry is then reduced to the 2 mm ideal symmetry.
drawn, as shown in Fig. 5, which were preliminary presented by the authors in a previous study [20]. Furthermore, although, Heino[31] and Heino et al. [32], and Lee et al. [14,15,33] have proven the formation of various carbides or other secondary phases, like R-phase, in these steels, none of these were identified in the present study. 3.2. Mechanical behavior The Vickers hardness curves of S31254 and S32654 steels were drawn, as shown in Fig. 6. Comparing the curves in Fig. 6A with those in Fig. 6B, it is shown that both steel grades exhibit similar hardness evolution at each temperature, with the hardness curves of S32654 shifted to higher values than those of S31254. Coarsening of precipitates over a critical size and dispersion occurs at shorter ageing times at 950 1C, where precipitation kinetics is faster. On the contrary, at 650 1C the hardness values seem to keep increasing even following 3000 h of ageing, implying that the peak on the curve might have not been reached as yet. Yield stress (s0.2%) and ultimate tensile stress (UTS) values of the AR specimens are lower than the nominal yield stress and UTS values, respectively. As shown in Fig. 7A, the yield stress and the UTS have either maintained (UTS values at 650 1C) or increased their values in all experimental conditions, compared to the AR values. The highest value was measured at 750 1C following 240 h of ageing. In addition, the elongation (e%) values decrease to the ageing time, while the AR elongation value is significantly higher than the nominal value (Fig. 7B). As an exception, no statistical difference was found between the elongation values measured at 950 1C, for either 120 h or 240 h. Fractography via SEM, revealed a combination of both ductile and brittle fracture in all specimens studied, with a more brittle behavior taking place following ageing at higher ageing temperatures and times. The modulus of toughness (UT) was calculated as the total area under the stress–strain curve [35]. Calculating the changes in UT, caused by precipitation, it is possible to evaluate the overall behavior of toughness versus ageing temperature and time. At any temperature, comparing the values obtained between 24 h and 240 h of ageing, the toughness decreases (Fig. 7C). However, no statistically significant differences were found between the 120 h of ageing and 24 h or 240 h.
4. Discussion 4.1. Precipitation of secondary phases in austenite The ageing temperature and time of the formation of the precipitates seem to be strictly dependent on the elemental composition of the steels. Even slight changes in the elemental composition of these steels are often the reason for somewhat different results between the various studies, on the same steel grades [14]. Small changes in the composition in Mo ( 71%) [31],
and even smaller in N ( 70.1%) [2,10,36] are sufficient to promote or retard the formation of different secondary phases in the same ageing times and temperatures, or similar secondary phases in different ageing times and temperatures. As shown in Fig. 5, precipitation of s-phase occurs at shorter ageing times in the most heavily alloyed with Cr and Mo S32654, when compared to S31254. In terms of ageing time, the first intermetallic phase that forms in both the steel grades, is the Laves phase, during ageing at either 650 1C or 750 1C, whilst at 950 1C Laves phase and s-phase nucleate almost at the same time in both steel grades, which is in agreement to previous studies [16,31]. Nevertheless, as Laves phase is a metastable phase, following ageing at 950 1C for 240 h very few Laves phase precipitates were observed. At the same time, the volume fraction of s-phase was increased, forming coarser particles, which implies either the transformation of Laves phase into s-phase, or the dissolution of Laves phase into the austenitic matrix followed by the formation of s-phase, or even both. Precipitation kinetics is much more sluggish at 850 1C, 750 1C and even more at 650 1C. Once Laves phase is formed, the same phase transformation is expected to occur at these temperatures, too, but at longer ageing times. It is believed that the dissolution of Laves phase is most possible to occur, as the relation between some of the crystallographic directions and planes of the hexagonal crystal system (Laves phase) and the fcc system of austenite is stronger than that of the tetragonal system of s-phase and the fcc system of austenite. At all the heat treatments performed in both steel grades the b-Cr2N nitride was always found adjacent to Laves phase particles and, given the fact that the Laves phase precipitates nucleate prior to the b-Cr2N nitride, the b-Cr2N nitride formation seems to be favored by the presence of Laves phase. Diffusion of Mo into austenite and precipitation of Laves phase causes the depletion of Mo in the surrounding matrix and promotes the formation of Cr-rich phases, such as s-phase and the b-Cr2N, turning the austenitic matrix richer in Ni as shown in Table 2. This could be an explanation for the formation of b-Cr2N particles adjacent to Laves phase particles (Fig. 2B). The b-Cr2N nitrides often form a continuous network of alternating precipitates along grain boundaries, which was also reported in previous studies [31]. Moreover, the initiation of the b-Cr2N nitride formation takes place at shorter ageing times in S32654 than in S31254, because of the higher content in N (Fig. 5). The cellular type of b-Cr2N was not observed in this study [1,33]. The b-Cr2N nitrides are more plate-like when observed intragranularly or adjacent to other types of precipitates (Fig. 2B) and more longitudinal when found single in the austenitic matrix and this is due to the faster diffusion of Cr in the direction of [0001]b:[111]g [19]. In the present study only a few longitudinal b-Cr2N precipitates were found alone in the austenitic matrix with such a shape. A few precipitates of w-phase were also found in specimens aged at 850 1C or 950 1C, a temperature range which is not very common for this phase transformation in stainless steels [29,37].
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Fig. 5. TTP diagrams for (A) S31254 and (B) S32654 [20].
1000 h of ageing at 850 1C, some w-phase precipitates were detected around the predominant s-phase precipitates. 4.2. Correlation of precipitation with mechanical behavior
Fig. 4. (A) Bright field TEM micrograph of a Laves phase precipitate observed following ageing of 1000 h at 850 1C in S31254 steel and (B) electron diffraction pattern recorded along [0001]L//[001]g from the particle of Fig. 4A, where (200)g: (2¯110)L and (020)g: [011¯0]L.
This phase has similar morphology to s-phase and it is often hard to distinguish them without the use of electron diffraction techniques. In addition, minimum or no defects were observed in the precipitates of w-phase, which is unusual to other types of stainless steels, such as duplex stainless steels [29], where w-phase is usually observed at lower temperatures with a lot of defects in its crystal structure. The volume fraction of w-phase does not increase as much as that of s-phase, but still, after
Within the initial ageing hours at either 850 1C or 950 1C the size of the precipitates remains at submicron values and precipitates form in a uniform dispersion, which results in an increase of hardness values. At 850 1C, hardness values increase up to a maximum following ageing for 120 h (Fig. 6), in both steel grades, where the combination of size and dispersion of Laves phase, s-phase and also b-Cr2N nitrides is optimized favoring hardness. In addition to precipitation hardening, another factor that assists the increase of hardness is hardening via second phase formation, since the volume fraction of the hard s-phase has been significantly increased and the size of the precipitates reach up to 10 mm. Following 120 h of ageing, partial dissolution or transformation of Laves phase to s-phase and coarsening of s-phase precipitates results to an overall decrease in hardness. The same behavior was observed during ageing at 950 1C, but coarsening of both Laves phase and s-phase and dissolution of Laves phase occurs at shorter time intervals. Moreover, increased amounts of Mo and Cr in S32654 lead to the formation of higher volume fraction of s-phase and Laves phase, than those formed in S31254. The decrease in hardness at 950 1C between 96 h and 120 h is statistically significant (Po0.001) and could be attributed to the change in hardness mechanism. It is the time period where the volume fraction of Laves phase decreases intensely and Mo dissolves to the austenitic matrix. Following that time interval, s-phase is primarily responsible for the further evolution of hardness.
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Fig. 6. Hardness evolution curves (HV) for (A) S31254 and (B) S32654.
As precipitation kinetics is slower and also different at 750 1C, during the first 24 h of ageing no s-phase was formed in S31254 and a few particles of s-phase were formed in S32654, with Laves phase being the predominant phase, but still forming in dispersed precipitates primarily along the grain boundaries. These precipitates may cause a small amount of precipitation hardening, but possible stress relief at 750 1C overcomes precipitation hardening and causes a decrease in overall hardness values. Also, as shown in Fig. 2A, mechanical twins were observed via TEM examination, caused by the growth of s-phase, which indicates that s-phase is hard enough to generate stresses at the s/g interface causing twinning of austenite. On the contrary, precipitation hardening during the initial ageing hours at 850 1C or 950 1C prevailed, resulting in a very fast increase of hardness values. The combination of different mechanisms that result into an overall decrease of hardness in both steel grades is better shown schematically by the ageing curves at 650 1C (Fig. 6). Up to 240 h of ageing hardness values lie below the AR value (P o0.001) while no statistically significant difference was found between the AR value and those measured at 500 h of ageing. This implies that original hardness value is achieved following 500 h of ageing. Precipitation kinetics at 650 1C is very slow with Laves phase being the predominant phase identified. Laves phase precipitates nucleate and coarsen at a very slow rate and hardness values increase over the AR values, only following 500 h of ageing in both steel grades. Few s-phase precipitates and b-Cr2N nitrides were observed at the specimens exposed for 3000 h of ageing, which is the longest ageing time for the present study. This substantially means that only Laves phase is involved in the evolution of mechanical behavior at 650 1C.
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As shown in Fig. 7A both UTS and yield stress values have been increased under all heat treatment conditions when compared to the AR values. The behavior at 650 1C is an exception, where both yield stress and UTS have either maintained or decreased their values, compared to the AR state, because no precipitation of the s-phase was observed and the volume fraction of Laves phase precipitates was minor. The overall strength of the steel was thus maintained or decreased because of possible stress relief that could take place during the initial ageing hours. Tensile behavior is similar to hardness behavior, meaning that when hardness values increase, UTS and yield stress values increase and vice versa. The optimum combination of Laves phase and s-phase dispersion and particle size leads to the maximum values of UTS and yield stress which correspond to the intermediate, for this study, at temperatures of 750 1C or 850 1C. Following heat treatments at 650 1C, slow precipitation kinetics of Laves phase and the absence of other types of secondary phases result in lower UTS and yield stress values, compared to the other ageing temperatures (Fig. 7A). Increased values of UTS and yield stress have turned the steel less ductile, as can be concluded by Fig. 7B. At 750 1C or 850 1C, the coarse s-phase particles and a uniform dispersion of Laves phase particles are responsible for a reduction in ductility values, compared to the AR value. Following ageing at 850 1C for 240 h, the elongation of S31254 is reduced by more than 65% and this is primarily attributed to the large, brittle precipitates of s-phase, but also to the precipitates of Laves phase and to a lesser degree to that of the dispersed b-Cr2N nitrides. The only experimental condition where elongation maintained the AR value is ageing at 650 1C for 24 h, because the small volume fraction of Laves phase precipitates has no major effect on ductility. However, considering the minimum nominal value for ductility of S31254 (Fig. 7B), it could be stated that the steel essentially maintains its elongation requirements following ageing at 650 1C for up to 240 h. Furthermore, the change in the modulus of toughness (UT), which is the amount of the work per unit volume versus ageing time and temperature, was also studied. The UT values are calculated by the total area under the stress–strain curve and should not be confused or compared to the impact of toughness values, which have been measured in previous studies [38,39], because it includes not only the energy absorbed during fracture but also the energy absorbed during plastic deformation. The evolution of UT versus ageing time and temperature (Fig. 7C) has similar behavior to the evolution of elongation (Fig. 7B), with all UT values calculated below the AR value of the steel in all experimental conditions. As an exception, toughness is essentially maintained following ageing at 650 1C for 24 h, as no precipitations were observed and no statistical significance was calculated for this heat treatment (P 4 0.05). Laves phase and s-phase precipitates dramatically affected the UT values, decreasing the values more than 30% of the AR values, following ageing at 850 1C for 240 h. This decrement was reasonable as elongation was found to decrease and so did the total area below the stress–strain curve, even if UTS was found to increase. As an overall comment for the present study it can be derived that Laves phase and/or s-phase are primarily responsible for all the changes in the mechanical properties of the aged superaustenitic stainless steels, while minor effect could be attributed to either w-phase or the b-Cr2N nitride. Following ageing above 750 1C both steel grades suffer significant changes in the mechanical properties, deteriorating their ductility but improving their strength, while it is much safer to expose these steel grades below 650 1C, where precipitation kinetics is much slower and mechanical properties are substantially improved, compared to higher temperatures.
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Fig. 7. Evolution of (A) yield stress s0.2% (MPa) and UTS (MPa), (B) elongation e (%) and (C) modulus of toughness UT (TPa) versus ageing time t (h) for S31254 (nom.: nominal value, AR: as measured in the as received condition).
5. Conclusions
At least four different types of precipitates, s-phase, w-phase,
Laves phase and b-Cr2N nitride, were identified following ageing of S32654 and S31254 within the temperature range of 650– 950 1C. At all temperatures, Laves phase is the first secondary phase forming, followed by s-phase and then at longer ageing times w-phase and the b-Cr2N nitride also form. Secondary phases form at earlier ageing times in the most heavily alloyed S32654. Full transformation of Laves phase to s-phase and/or its dissolution in the austenitic matrix takes place for time intervals over 240 h at an ageing temperature of 950 1C. Precipitation of primarily Laves phase and s-phase resulted into an increase of the hardness values at all temperatures in both the steel grades. Secondary phase precipitation resulted into an increase of both yield stress and UTS while elongation and toughness were decreased in both steel grades.
Appendix A. Supplementary materials Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.msea.2012.10.066.
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