Volumetric changes associated with B2-(Ni,Fe)Al dissolution in an Al-alloyed ferritic steel

Materials and Design 111 (2016) 640–645

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Volumetric changes associated with B2-(Ni,Fe)Al dissolution in an Al-alloyed ferritic steel Reza Rahimi a, Péter Pekker b, Horst Biermann c, Olena Volkova a, Bruno C. De Cooman d, Javad Mola a,⁎ a

Institute of Iron and Steel Technology, Technische Universität Bergakademie Freiberg, Germany Institute of Metal Forming Science and Nanotechnology, University of Miskolc, Hungary c Institute of Materials Engineering, Technische Universität Bergakademie Freiberg, Germany d Graduate Institute of Ferrous Technology, Pohang University of Science and Technology, Pohang, Gyeongbuk 790-784, South Korea b

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• Anomalous length changes during continuous dilatometry heating of an Alalloyed ferritic stainless steel were studied. • The anomaly in the coefficient of thermal expansion was correlated with the precipitation and dissolution of B2(Ni,Fe)Al intermetallics. • Increase in the Al and Ni content of the ferritic matrix due to the dissolution of B2 intermetallics was responsible for the high apparent coefficients of thermal expansion.

a r t i c l e

i n f o

Article history: Received 23 May 2016 Received in revised form 20 August 2016 Accepted 9 September 2016 Available online 10 September 2016 Keywords: Al-alloyed steels Stainless steels Dilatometry Coefficient of thermal expansion B2-intermetallics

a b s t r a c t An Al-alloyed Fe–Cr–Ni–Al–Mn–C ferritic stainless steel exhibited a high apparent coefficient of thermal expansion during continuous heating in the temperature range of 800–1050 °C. The anomaly was ascribed to the matrix lattice expansion associated with the dissolution of (Ni,Fe)Al precipitates with a B2-CsCl crystal structure. The expansion was justified on the basis of the Al and Ni enrichment of the ferritic matrix as the precipitates dissolved. © 2016 Elsevier Ltd. All rights reserved.

1. Introduction Steels containing up to about 10 mass-% Al have been designed for applications requiring a combination of high strength and ductility [1, 2]. The density reduction caused by Al addition makes such steels ⁎ Corresponding author. E-mail address: [email protected] (J. Mola).

http://dx.doi.org/10.1016/j.matdes.2016.09.033 0264-1275/© 2016 Elsevier Ltd. All rights reserved.

attractive for a variety of applications particularly in the automobile industry where steel design concepts are driven by reduced fuel consumptions and exhaust emissions [3–6]. In cases where an austenitic rather than a duplex or a ferritic microstructure is desirable, the high ferrite potential of Al is compensated by the addition of a high carbon concentration [7]. The simultaneous presence of high concentrations of Al and C in high Mn steels in turn favors the formation of (Fe,Mn)3AlCx kappa carbides in the austenite [8]. Kappa carbide is also

R. Rahimi et al. / Materials and Design 111 (2016) 640–645

the dominant type of carbide in ferritic steels containing Al, Mn, and C [2,9]. The ferrite phase of Al-alloyed high Mn duplex steels where the carbon partitions into the austenite, on the other hand, is prone to the formation of ordered intermetallic compounds most notably those of the type B2-FeAl and D03-Fe3Al [10]. In particular, the formation of B2-intermetallic domains in ferrite, favored by the Ni addition [11], has been associated with promising mechanical properties. In Al-alloyed ferritic steels containing high Ni contents, Fe in the B2FeAl is partially replaced by Ni [11]. The formation of B2-(Ni,Fe)Al intermetallic precipitates in Fe–10Cr–10Ni–3.4Mo–(3− 10)Al-base ferritic steels (values in mass-%) and its influence on the mechanical properties have been explored in many recent investigations [12–18]. In contrast to B2-intermetallics in high Mn steels which appears as domains in the ferrite hardly identifiable without diffraction analysis in transmission electron microscope (TEM) [8], B2-phase in the latter alloy system and in high Cr steels alloyed with Ni and Al [19–22] appears as discrete precipitates in the ferritic matrix. This enables their identification even by scanning electron microscopy. In certain commercial steels, for instance the stainless steel grade PH13-8 Mo, B2-intermetallics serve as strengthening second phase particles [23]. The dissolution and precipitation of B2-intermetallics have profound consequences for the mechanical properties of Al-alloyed steels [17,24]. Therefore, devising straightforward methods of studying the preceding reactions will eliminate the need for time-consuming microstructural examinations otherwise necessary for this purpose. Non-isothermal dilatometry is a powerful and yet quick method of analyzing solid-state phase transformation kinetics which relies on the identification of inflection points in dilatation data obtained at various heating rates [25]. This method has been used for instance for the analysis of tempering reactions in martensitic steels [25–27], dissolution of cementite in austenitic steels [28], and most commonly for the identification of critical temperatures such as A1, A3, and Acm at different heating and cooling rates [29,30]. The present paper demonstrates a dilatometry-based method of studying the dissolution and precipitation of B2-intermetallics in an Al-alloyed stainless steel.

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3. Results and discussions The microstructure after homogenization treatment is shown in Fig. 1. The coarse precipitates readily visible in the optical micrograph of Fig. 1(a) have been shown elsewhere to be of the type M7C3, where M denotes Cr, Fe, and substitutional alloying elements [31]. The SEM micrograph in Fig. 1(b) reveals the presence of fine precipitates in the matrix of Fig. 1(a). Supplemental TEM examinations presented in Fig. 1(c) confirmed a bimodal distribution of coarse (~ 200 nm) and fine (b20 nm) precipitates in the matrix. Selected area electron diffraction (SAED) patterns of a coarse precipitate and a region of the matrix containing only a small fraction of fine precipitates are shown in Fig. 1(d) and 1(e), respectively. Both SAED patterns include reflections pertaining to the ordered B2-intermetallic phase with a cubic CsCltype crystal structure. However, due to the presence of a lower volume fraction of B2-intermetallics, the intensity of their characteristic superlattice reflections is weaker in Fig. 1(e). The orientation relationship between the B2-intermetallics and the ferritic matrix is of the – – type cube-on-cube, namely (011)B2 // (011)α and [011]B2 // [011]α. Quantitative energy dispersive spectroscopy analysis (EDS) in TEM revealed that the B2-intermetallics were enriched with Al and Ni (see Table 2).

(a)

(b) M 7 C3

30 µm

1 µm

(c)

2. Methods

d

Al-added steel with the chemical composition shown in Table 1 was cast in a vacuum induction melting furnace. The cast ingot was homogenization annealed at 1200 °C for 1 h. Subsequent blow of N2 gas yielded an average cooling rate of 2.5 °C/s in the temperature range 1200–600 °C. The microstructure was subsequently examined using a Neophot 30 type light optical microscope, a Zeiss Ultra 55 type field emission scanning electron microscope (FE-SEM), and an FEI Tecnai G2 type transmission electron microscope (TEM) operated at an acceleration voltage of 200 kV. Samples for TEM examinations were mechanically polished to a thickness of b100 μm and then twin-jet polished at room temperature (RT). X-ray diffraction (XRD) measurements with Cu Kα radiation were performed in a Seifert-FPM URD-6 diffractometer. Dilatometry studies using 3.5 × 3.5 × 10 mm3 samples were done in a BÄHR-DIL805 dilatometer. Heating and soaking steps were done in vacuum before cooling with argon. Heating and cooling rates in dilatometry cycles were 50 °C/s and 20 °C/s, respectively. In-situ high temperature XRD measurements with Cu Kα radiation were conducted in a Bruker D8 diffractometer equipped with an Anton Paar heating stage. To minimize the oxidation, the specimen chamber was evacuated and backfilled with Ar.

Table 1 Chemical composition of the Al-added steel.

e

0.3 µm (d)

111B2

011B2

100B2

0 1 Z.A.

Element

C

Al

Mn

Cr

Ni

Si

Fe + others

mass-% at.%

0.46 1.93

7.1 13.3

6.2 5.7

17.1 16.6

8.7 7.5

0.4 0.7

Balance Balance

(e)

011α,B2

200 α,B2 100B2

0 1 Z. A.

Fig. 1. (a) Optical micrograph showing an overview of the homogenized microstructure. (b) SEM and (c) TEM micrographs of the matrix in (a). (d–e) SAED patterns taken from the regions marked in (c).

R. Rahimi et al. / Materials and Design 111 (2016) 640–645

Fig. 2 shows the first derivative of relative length change obtained during continuous dilatometry heating of the alloy. The derivation range was approximately 30 °C. In other words, relative length change data 15 °C below and 15 °C above the temperature of interest were used to construct the derivative curve. The derivative curve represents the apparent coefficient of thermal expansion (CTEa). For comparison, CTEa changes during continuous heating of a fully ferritic non-transformable Fe–18%Cr alloy are also included in Fig. 2. At temperatures below approximately 500 °C, CTEa values for both alloys remain comparable. Abrupt CTEa variations around 514 °C and 650 °C mark Curie temperatures (Tc) of the ferritic matrix and that of the reference Fe–18%Cr alloy, respectively [32]. Above Tc, CTEa of the experimental alloy remains above but almost parallel to that of the reference Fe–18%Cr alloy. At temperatures above approximately 800 °C, CTEa increases dramatically and diverges quickly from the CTEa curve for the reference Fe–18%Cr alloy. At 995 °C, CTEa is almost 50% higher than that of the reference ferritic alloy (29 × 10−6 °C−1 vs 19 × 10−6 °C−1). Above 995 °C, CTEa begins to decrease to a local minimum of 24 × 10−6 °C−1 at 1050 °C before increasing again at higher temperatures. Kinetics of the process responsible for the CTEa anomaly was investigated by dilatometry heating at different rates. As shown in Fig. 3(a), as the heating rate changes in the range 0.1–50 °C/s, the temperature associated with maximum CTEa and the subsequent CTEa decrease remains unaffected. Furthermore, Fig. 3(b) reveals that the phenomenon responsible for the observed CTEa variation between 800 °C and 1050 °C is highly reversible such that the CTEa curves during dilatometry heating and cooling almost overlap. In order to study the origin of abnormally high CTEa values above 800 °C and the subsequent CTEa decrease between 995 °C and 1050 °C, small specimens were heat treated at different temperatures and quenched in water. The applied heating rate and soaking time for the latter furnace heat treatments were 0.3 °C/s and 10 min, respectively. Water quenching aimed at freezing the high temperature microstructures. Fig. 4(a)–(c) show SEM micrographs of the specimens quenched from the temperatures marked by vertical arrows on the CTEa curve of Fig. 2. At 800 °C, the ferritic matrix contains a high fraction of B2-intermetallics. After quenching from 975 °C, slightly below the peak CTEa, the fraction of B2-intermetallics is almost half of that at 800 °C. After heat treatment at 1100 °C which is above the local CTEa minimum, B2intermetallics were absent near the surface of the water-quenched specimen where the cooling rate was highest. The results confirm that the abnormal CTEa values between 800 °C and 1050 °C are governed by the dissolution of B2-intermetallics. Completion of B2-intermetallics dissolution at 1050 °C is in fair agreement with the dissolution temperature of B2-intermetallics in stainless steels alloyed with Ni and Al [21]. In binary Fe–Al alloys with Al contents of 36–49 at.%, an ordered B2FeAl structure is stable at temperatures up to about 1200 °C [33]. B2FeAl phase in such binary steels is capable of generating a high concentration of thermal vacancies [34,35]. The generation of vacancies may be thought of as the migration of atoms from lattice sites to free surfaces. An increase in the concentration of thermally-induced vacancies would be associated with an expansion in the sample volume which can be detected by length change measurements in a dilatometer [36, 37]. Comparison of dilatometric length changes with the lattice parameters obtained from high temperature diffraction measurements will then allow to quantify the concentration of thermal vacancies [35]. Due to the presence of B2-intermetallics in the present alloy, the high CTEa values in the temperature range of 800–1050 °C may at first glance

Fe-18mass-%Cr Al-alloyed steel

0.30

Apparent CTE (10-4 °C-1)

642

Fig. 4(b) Fig. 4(c)

Heating Section

0.25 Fig. 4(a)

End of CTE a anomaly

0.20 TC 0.15

0.10

+50

°C/s

200

400

600

800

1000

Temperature (°C) Fig. 2. CTEa values during continuous heating of the experimental alloy and a reference Fe − 18%Cr ferritic stainless steel. The microstructures corresponding to the marked temperatures are shown in Fig. 4.

appear to be due to the formation of thermal vacancies in B2-intermetallics. In this case, the possible contribution of thermal vacancy formation in B2-intermetallics to the overall dilatation of the alloy would be proportional to their volume fraction. Assuming that the entire Al in the present alloy participates in the (Ni,Fe)Al formation, the maximum fraction of B2-(Ni,Fe)Al would be only about 27 vol.%. Therefore, even if B2-(Ni,Fe)Al carried a high concentration of thermal vacancies, the observed decrease in their volume fraction would lead to a contraction and a CTEa decrease. Therefore, thermal vacancy formation in B2(Ni,Fe)Al whose volume fraction continuously decreases cannot justify the observed high CTEa values. The CTEa anomaly was further studied by in-situ XRD measurements at temperatures up to 1050 °C. The specimen used for this purpose was heat treated at 1200 °C for 15 min followed by quenching in a mixture of water and 10 mass-%NaCl (brine-quenched). Brine quenching was intended to freeze the high temperature microstructure and minimize the fraction of B2-(Ni,Fe)Al which might form in the cooling step. The specimen was then crushed into powder and was used for in-situ XRD measurements. The first and last XRD measurements at RT were done over a scan range of 30° b 2θ b 140°. To minimize the specimen exposure to high temperatures and the potential risk of oxidation, XRD measurements at other temperatures were limited to a small range around 110α and 211α diffraction peaks. The average heating and cooling rates were +0.16 °C/s and −0.3 °C/s, respectively. The interplanar spacing of 211α planes (d211) and the corresponding CTE values for ferrite (CTEbcc) during heating and cooling are plotted in Fig. 5(a)–(c). Due to the high cooling rate in the prior brine quenching step and the low fraction of B2-(Ni,Fe)Al prior to heating, the interplanar spacings during heating differ from those during the slow cooling. The RT lattice parameters before and after in-situ XRD measurements were determined by the application of the Nelson-Riley function [38] to be 0.28893 (±0.00003) nm and 0.28807 (±0.00003) nm, respectively. This is due to the difference in the volume fraction of B2-(Ni,Fe)Al precipitates which in turn influences the chemical composition and the lattice parameter of the ferritic matrix. As exemplified in Fig. 2 with an Fe-18%Cr steel, CTEa of ferritic steels gradually increases with temperature. However, CTEbcc in the heating

Table 2 Chemical composition of a region of the matrix and B2-intermetallics based on the TEM-EDS analysis. Element

Matrix B2-intermetallic

Al

Mn

Cr

Ni

Fe + others

mass-%

at.%

mass-%

at.%

mass-%

at.%

mass-%

at.%

3.5 19.0

7.0 33.3

7.6 6.9

7.3 5.9

15.3 2.6

15.7 2.4

2.6 53.4

2.3 43.0

Balance Balance

R. Rahimi et al. / Materials and Design 111 (2016) 640–645

0.30

0.30

0.28

0.25

0.26

0.24

900

920

940

960

980

1000 1020

0.20 0.1 °C/s 0.5 °C/s 1 °C/s 5 °C/s 10 °C/s 25 °C/s 50 °C/s

0.15 0.10 200

400

600

800

1000

0.40

Heating Cooling 0.30

°C /s

0.32

Apparent CTE (10-4.°C-1)

Apparent CTE (10-4.°C-1)

0.35

0.28

0

(b) 0.34

-2

(a)

643

0.26

0.24

0.22 850

C/s

0°

+5 900

Temperature (°C)

950

1000

1050

1100

Temperature (°C)

Fig. 3. (a) CTEa values as a function of temperature at various heating rates. The inset shows a magnified view of curves around the CTEa maximum. (b) CTEa changes in the vicinity of the CTEa maximum during heating and cooling.

segment of in-situ XRD (Fig. 5(a)) undergoes a decrease in the temperature range marked I in Fig. 5(b). Due to the high supersaturation of ferrite in the brine-quenched specimen with respect to Ni and Al, release of the Al and Ni supersaturation by the formation of B2-(Ni,Fe)Al is expected to occur in the subsequent heating step. A decrease in the solute content of both elements, particularly Al, decreases the lattice parameter of ferritic steels [39,40]. For instance, the lattice parameter of an Fe-5 mass-%Al binary steel is about 0.5% higher than that of pure iron (0.2880 nm vs 0.2866 nm) [40]. Accordingly, the CTEbcc decrease marked I must be related to the associated Al and Ni depletion of ferrite due to the B2-(Ni,Fe)Al formation. In the temperature range marked II in Fig. 5(b), CTEbcc increases dramatically. The high rate of increase in the lattice parameter of ferrite may be justified by its enrichment with respect to Al and Ni as the dissolution of B2-(Ni,Fe)Al precipitates proceeds. This confirms that the high dilatometry-based CTEa values between 800 °C and 1050 °C are in fact due to the lattice expansion of the ferritic matrix as B2-intermetallics dissolve. The slight CTEbcc decrease observed in the temperature range marked III in Fig. 5(b) is expected to reflect a decrease in the rate of Al and Ni enrichment of ferrite at the latest stage of B2-(Ni,Fe)Al dissolution. The reduction in the dilatometry CTEa values at temperatures above approximately 995 °C could then be related to a reduction in the fraction and the rate of B2-(Ni,Fe)Al dissolution. Above 1050 °C, where the B2-(Ni,Fe)Al dissolution and the Al and Ni enrichment of ferrite no longer play a role, CTEa variations are primarily governed by the lattice expansion and the possible vacancy formation in ferrite. Quantitative metallography indicated that the fraction of carbides after brine quenching from 1250 °C was 6.5 vol.%, which is slightly lower than the carbide fraction in the homogenized condition of 7.9 vol.%. Therefore, the possibility of a slight increase in the solute C

(a)

1 µm

800 °C (b)

1 µm

and Cr concentration of ferrite due to the gradual dissolution of carbides cannot be excluded. Consistent with the above analysis, the formation of B2-intermetallics during the reverse process, namely aging of as-quenched ferritic stainless steels containing Ni and Al, has been found to reduce the ferrite lattice parameter [21]. In the present case too, the Al and Ni depletion of ferrite during in-situ cooling from 1050 °C resulted in a high rate of ferrite lattice contraction which manifested itself as high CTEbcc values in the cooling step (Fig. 5(c)). At temperatures below approximately 600 °C, CTEbcc values decreased to levels typical of ferritic steels. According to the preceding discussions, the higher lattice parameter of ferrite in the brine-quenched condition compared to that in the homogenized condition is due to the higher solute Al and Ni concentration of ferrite. According to in-situ XRD measurements, the dilatometric length changes correlate well with the changes in the lattice parameter of ferrite. Therefore, the CTEbcc decrease observed in the in-situ XRD results (temperature range marked I in Fig. 5(b)) must also occur during dilatometry heating of brine-quenched specimens. In Fig. 6(a), the dilatometry CTEa plot subsequent to brine-quenching is compared with that in the as-homogenized condition. The former shows a clear CTEa decrease in the vicinity of 700 °C where the CTEbcc decrease occurred too. The coincidence of the signals from dilatometry and in-situ XRD indicates that dilatometry is well suited for the study of the B2-(Ni,Fe)Al precipitation and dissolution. Due to the higher solute content of ferrite in the brinequenched specimen, it shows a lower Tc compared to the as-homogenized specimen (514 °C vs 483 °C). The present case of ferritic matrix dilatation during B2-(Ni,Fe)Al dissolution is analogous to the austenite dilatation during dissolution of cementite in hypereutectoid steels [28]. In the latter case, in spite of the higher atomic volume of cementite compared to that of the austenite,

975 °C (c)

1100 °C

1 µm

Fig. 4. Microstructure of Al-alloyed steels after annealing of the homogenized ingot for 10 min at marked temperatures. Samples were electro-etched with pure HNO3.

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R. Rahimi et al. / Materials and Design 111 (2016) 640–645

(a)

Heating (XRD) Cooling (XRD)

0.1205

d 211 (nm)

0.1200 0.1195 /s

0.1190

°C

6

1 0.

+

/s

0.1185 .3 -0

°C

0.1180 0

200

400

600

800

1000 1200

Temperature (°C)

(b)

(c)

0.30 0.28

XRD_Heating

0.30 0.28

III

XRD_Cooling

0.26

0.24

CTE bcc (10-4.C-1)

CTE bcc (10-4.C-1)

0.26

II

0.22 0.20 0.18

I

0.16 0.14

0.24 0.22 0.20 0.18 0.16 0.14

0.12

0.12 500

600

700

800

900

1000

0

200

Temperature (°C)

400

600

800

1000

Temperature (°C)

Fig. 5. (a) Interplanar spacing of 211α planes during in-situ heating and cooling of a brine-quenched specimen; corresponding CTEbcc values in (b) heating and (c) cooling segments calculated based on the interplanar spacings in (a).

the simultaneous carbon enrichment of austenite during the cementite dissolution leads to CTEa values higher than the values characteristic of single-phase austenitic steels. In other words, if the dissolution is allowed to occur isothermally, it will lead to a net expansion. In the exemplified derivative dilatometry curve of an Fe−6Mn−1.5C (mass-%) austenitic steel (Fig. 6(b)), the formation of cementite in the austenite leads to initially low CTEa values and its progressive dissolution at higher temperatures leads to CTEa values higher than those of singlephase austenitic steels, which are commonly in the range of (20–

0.20

B2 dissolution Tc 0.15 800

γ only

0.35

Heating Segment

0.30 0.25

o

0 +2

C/

M 3C

s

precipitation M 3C dissolution

0.20 0.15

B2 precipitation 600

0.40 Austenitic Fe-6Mn-1.5C steel

Heating Section +50 °C/s

Temperature (°C)

In summary, unusually high CTEa values of an Fe–Cr–Ni–Al–Mn–C stainless steel in the temperature range of 800–1050 °C was rationalized by the dissolution of existing B2-intermetallics of the type (Ni,Fe)Al. Although Al-rich B2-intermetallics are known to be capable of generating

(b)

Brine-quenched from 1200 oC As-homogenized

0.25

400

4. Conclusions

Apparent CTE (10-4 °C-1)

Apparent CTE (10-4 °C-1)

(a) 0.30

25) × 10−6 °C−1 [32,41]. It is only after completion of cementite dissolution that the CTEa recovers to the expected level.

1000

200

400

600

800

1000

Temperature (°C)

Fig. 6. (a) Comparison of CTEa values during continuous heating of as-homogenized and brine-quenched specimens; (b) CTEa values during continuous heating of an austenitic Fe−6Mn−1.5C (mass-%) steel.

R. Rahimi et al. / Materials and Design 111 (2016) 640–645

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