Structural mechanism of reverse α → γ transformation and new functional properties of Fe-Ni austenitic alloys

V.V. Sagaradze, I.G. Kabanova, N.V. Kataeva, and M.F. Klyukina

Institute of Metal Physics, Ural Branch, Russian Academy of Science,

Ekaterinburg 620990, Russia

E-mail: vsagaradze@imp.uran.ru

Keywords: Fe-Ni austenitic alloys, cyclic martensitic transformation γ→ε,α→γ, ordering, nanocrystalline γ phase, strengthening, magnetic properties, Ni redistribution, thermal expansion coefficient

Abstract. The structure of the metastable austenite Fe–32%Ni alloy quenched for martensite and subjected to the α→γ transformation at slow heating (to 593–773 K) with the formation of a nanocrystalline austenite of different orientation was investigated. Electron diffraction analysis has revealed that a nickel-enriched nanocrystalline γ-phase acquires ordered L10 superstructure. Together with the γ-phase, there was detected the disperse ε-martensite with a hexagonal close-packed (hcp) lattice. Along with the realization of ordinary orientation relationships (ORs) of Kurdjumov-Sachs, between a matrix α-phase and an ordered γ-phase in the sample of Fe-32%Ni alloy after slow heating to 773 K there have been established Headley–Brooks ORs. Cyclic transformation promotes the improvement of many functional properties of steels (strength, coercive force, controlled thermal expansion coefficient, and others). Introduction

Nanocrystalline austenite in the bulk of any sample of Fe-Ni metastable austenite alloys can be obtained in result of multiplication of γ-phase orientations (to 576) in the process of the direct γ→α and reverse α→γ transformations [1-5]. Various phase transformations and new functional properties (magnetic-disperse hardening, formation of nanocrystalline austenite, “phase hardening”, the orientation of γ-phase habit planes, the specific features of the formation and inheritance of dislocation structures, the conditions of formation of a nickel-enriched austenite, and others) have been observed in these alloys [1-6]. The authors of [7-11] have pointed to that at particular conditions in Fe-Ni alloys there exists a possibility of the ordering according to the L10 type. The ordered phase Fe–50%Ni (tetrateinite) is often observed in Fe-Ni meteorites [7-9].

The present study is devoted to the electron diffraction analysis of the process of α→γ transformation at a slow heating in the alloy Fe–32%Ni with the aim to detect an ordering of the γ-phase according to the L10 type. The authors expected to detect intermediate ε martensite at the α→γ transformation with taking into account that at slow heating (0.01-0.2K/min) this transformation proceeds in several stages and its interval widens and shifts towards the region of higher temperatures: from 673 to 823 K [2]. We also investigated the changes in the physico-mechanical properties of Fe-Ni alloys after the γ → α → γ transformation. Experimental

An iron alloy with 32.5 wt % Ni and 0.01 wt % C (Fe-32Ni) and martensite point Мs = 180K was taken for investigation of the structure and phase transformations. Samples were quenched in water from 1373 K and cooled in liquid nitrogen for the formation of ~80% athermic lenticular α-martensite with a bcc lattice. Nanocrystalline austenite was formed during the α→γ transformation at a slow heating with a rate of 0.01 and 0.2 K/min from 293 K to different temperatures (593–773 K) within the critical interval. Transmission electron microscopy (TEM) analysis was carried out on foils of Fe-32Ni alloy using electron microscope JEM-200CX. The foils were prepared by

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mechanical grinding of samples to a thickness of ~100 µm with the subsequent electropolishing in the orthophosphoric acid with chrome anhydride dopes. An interpretation of diffraction patterns, an analysis of the reciprocal-lattice sections and projections of the nodes onto a diffraction plane, were carried out with using the computational program «CaRIne v. 3.1». Mechanical, magnetic, and thermal properties were investigated after γ→α→γ transformation. Results and discussions

Phase transformation α → γ.

As was reported in [1-5], during the α→γ transformation in the Fe-32Ni alloy at a slow heating with a rate of 0.2 K/min one can observe the formation of up to 24 orientations of the nanocrystalline austenitic plates with a thickness of ~ 5-10 nm (at initial stages of transformation) and an enhanced nickel concentration (to 40-50 wt %Ni). Such a multiplication of γ-phase orientations at the realization of Kurdjumov-Sachs orientation relationships (ORs) takes place in the case of elimination of the orientation-determining influence of a retained austenite. During slow heating, a “buffer” layer of low nickel content is formed along the boundaries of martensite crystals, which hampers the epitaxial nucleation of the γ-phase in the retained austenite as on a substrate [2-5]. The structure with differently oriented γ-phase nanocrystals that is typical of the Fe-32Ni alloy is shown in Figs.1a, 1b. Crystals of a nano-sized γ-phase (аγ = 3.585 Å) are arranged within the initial martensite plates (аα = 2.87 Å) (Fig. 1a). Thin γ-phase plates (Fig. 1b) have an inner banded structure and consist of separate fragments and, possibly, microtwins.

Fig.1. Structure (a) of α-martensite crystals and (b) of the disperse γ phase formed in result of α→ γ transformation at a slow heating with a rate of 0.2K/min to 753K in the alloy Fe-32Ni alloy.

The structural mechanism of formation of a nanocrystal γ-phase at slow heating is considered. Fig. 2 (a-c) shows the electron diffraction pattern and structure of dislocation part of a lenticular α-martensite crystal in Fe-32Ni alloy after slow heating at a temperature lower than Аs with a rate of 0.01K/min to 593K. In the bright-field and dark-field images (Fig. 2 a,b) inside the α crystal with the zone axis (ZA) [111]α, we can see thin crystals of ε-martensite of several orientations to which the reflections (010)ε, (002)ε, and (011)ε in the electron diffraction pattern correspond. In spite of the high energy of stacking faults in the Fe-32Ni alloy, the appearance of an intermediate ε-phase with hcp lattice (а=2.535 Å, с=4.132 Å, с/а=1,63) precedes the formation of the disperse austenite. Apparently, the occurrence of an intermediate ε-phase at the α→γ transformation has a more general character and is also typical of the iron alloys with high values of the energy of stacking faults, as was described in [12].

At the similar temperature of slow heating (593 K), in the area of twin midrib we observe a partial untwinning of α-martensite crystals. It happens discreetly with formation of thin plates of matrix orientation that cut transversely the initial α-twins (Fig. 3a). Since in the electron diffraction pattern (Fig. 3b) there are reflections not only from matrix and twin but also from ε-phase, we have grounds to believe that in the dark-field image in a combined (α+ε) reflection not only plates of “split” twin but also thin ε-crystals are shined.

1 µm

а

500 nm

b

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a b c d Fig.2. Structure (а) of dislocation part of a lenticular α-martensite crystal in Fe-32Ni alloy under slow heating with a rate of 0.01 K/min to 593K; b - electron diffraction pattern with reflections from α- and ε- phases; c – dark-field image taken in combined reflection (d): (01 1 )α+(010)ε+(002)ε+(011)ε.

а b Fig.3. Structure (а) in the area of twin midrib of an α-martensite crystal in the Fe-32Ni alloy under a

slow heating to 593K with a rate of 0.01K/min; (dark field in the combined reflection ( ε)111( + tw)110( α ); b - electron diffraction pattern with reflections from α- and ε- phases.

Slow heating of up to 673–723 K leads to the formation of a nano-sized γ-phase with fcc

lattice (Fig. 4 a-c). In dislocation area, the crystals of all 24 possible γ-orientations [2-5] are formed in accordance with the martensite orientation relationships (ORs) similar to Kurdjumov-Sachs relationships. In the area of partially remained midrib, the number of forming γ-phase orientations is limited (Fig. 4 a), since austenite crystals are expected to be simultaneously connected by martensite ORs with a matrix and a twin [3-5]. It should be noted that under such a slow heating in the Fe-32Ni alloy the forming γ-phase is enriched by nickel to 40-50% [5] and experiences the ordering of L10 type. Figure 4 shows the electron diffraction pattern with superstructure reflections of the (001)γ and (110)γ types (Fig. 4 c) and the dark-field image of nanocrystal layered austenite taken in the reflection (001)γ of the ordered γ-phase (Fig. 4 b). The L10 ordering in thin-layer γ-phase of the Fe-32Ni alloy is detected at a temperature of 723 K which is much higher than the calculated value of Тс = 603 K [8] and this requires an improvement of theoretical calculations. The electron diffraction pattern in Fig. 4c shows the reflections not only from the initial α-martensite and γ-austenite but also from the ε-phase.

200 nm

200 nm

[ ] [ ] [ ]εαα 110;223;410 ZAZAZAtw

α141

α141

tw

α334tw

α244

ε002

ε111

ε111α200

tw

α110

ε101

200 nm

( )ε001[ ]ε001

α110

α101

α101

[ ]α111ZA

010ε

002ε

α101

011ε

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Along with the realization of ordinary orientation relationships (ORs) of Kurdjumov-Sachs, between a matrix α phase and an ordered γ-phase in the sample of Fe–32%Ni alloy after slow

heating to 773 K there have been established Headley-Brooks ORs [13]: ( 011 )α ║ ( 111 )γ, ZA[110]α ║ZA[110]γ, (Fig. 4d).

At the final stage of the α→γ transformation, globular austenite is formed (Fig. 5) that absorbs the nickel-enriched γ-plates and the nickel-depleted martensite. In the nonuniform structure of the austenite, fringe diffraction contrast is seen (Fig. 5), which reflects the existence of alterations in the nickel concentration. The homogenization of the alloy in the austenitic state can be provided by a 1-h holding at 1100 Κ.

a b

c d Fig.4. Dark-field image of nano-sized austenite in Fe-32Ni alloy in (а) the reflection (111)γ and in (b) the superstructure reflection (001)γ; (c) electron diffraction pattern corresponding to Fig. b and (d) electron diffraction pattern corresponding to Headley–Brooks relationship. Treatment: α→ε,γ transformation under heating with (а-с) a rate of 0.01 K/min to 723 K and with (d) a rate of 0.2 K/min to 773 K.

Fig.5 Formation of globular austenite in Fe-32Ni alloy in the final stage of α→γ transformation under slow heating to 763K with a rate of 0.2K/min. Changes in physico-mechanical properties upon γ→α→γ transformation

The formation of differently oriented nanocrystalline lamellar austenite in the process of α→γ transformation upon slow heating can even lead to strengthening of α martensite, since these γ-crystals represent serious obstacles for moving dislocations. Thus, the formation of 15% nanocrystalline γ-phase inside the crystals of α-martensite upon heating to 770 Κ increases the yield

1 µm

200 nm 200 nm

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strength of the 32Ni steel from 670 to 1100 MPa (Fig. 6, curve 1). Note that in this case it is "uphill" hardening of: the martensite is strengthened in result of the formation of disperse austenite in it. However, the nucleation of austenite on the boundaries of martensite laths (as is observed in the 28Ni alloy with a structure of packet martensite) does not lead to a significant increase in the strength properties (Fig. 6, curve 2). At the final stage of the α→γ transformation, when globular austenite is formed, the Fe-Ni alloys exhibit sharp softening (Fig. 6).

Fig. 6. Variations in the yield stress σ0.2 (at 293K) and amount of the austenite formed upon α→γ transformation using slow heating at a rate of 0.3 K/min in (1) Fe-30Ni and (2) Fe-28Ni alloys depending on the heating temperature.

Fig. 7. Minor hysteresis loops in the 18Ni-9Co-5Mo-1Ti steel measured in a field H of (1-3) 32 and (4) 65 A/cm, respectively. Treatment: (1) quenching from 1273 K in water; (2) quenching + cold rolling to ε = 40%; (3, 4) quenching + rolling (to 40%) + partial α→γ transformation and aging (823 K, 2 h + 843 K, 1 h).

The formation of fine lamellar γ-crystals upon the α→γ transformation makes it possible to

"disintegrate" the domain structure of the α-martensite and change some magnetic properties of the alloys. In particular, the partial formation of disperse γ-phase in the process of α→γ transformation upon heating to 823-843 K in the aged textured sheet of a maraging steel 18Ni-9Co-5Mo-1Ti leads to an increase in the coercive force Hc and retention of a highly rectangular hysteresis loop in small fields [14]. Thus, the high-strength steel 18Ni-9Co-5Mo-1Ti (170 MPa) can be transformed with passing into the category of "semihard"-magnetic materials, which are suitable, in particular, for the use in high-speed hysteresis-reluctance motors.

The possibility of the formation of concentration-inhomogeneous austenite upon the α→γ transformation during slow heating (Fig. 5) can be used for controlling the thermal expansion coefficient (TEC) of Fe-Ni alloys within sufficiently wide limits [15]. Table 1 shows experimentally determined values of the TEC of the 32Ni alloy in the temperature range of 223-323 Κ after various heat treatments. A slow heating into the two-phase α + γ field to temperatures of 703-763 Κ causes partial α→γ transformation and realization of different degrees of Ni redistribution between the α- and γ- phases. A subsequent accelerated heating into the single-phase γ-field (873 K) from temperatures of 703-763 Κ fixes (in the austenite) different concentration separation in nickel. This is the cause for the appearance of different thermal properties in the 32Ni alloy. The TEC can be changed from 2.5×10-6 to 10.5×10-6 K-1 depending on the regime of heat treatment.

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Table 1. Average values of the thermal expansion coefficient (10-6 Κ-1) of the Fe-32Ni alloy in various temperature ranges after various heat treatments

№ Heat treatment Temperature range, Κ

223-323 523-623

1 Quenching from 1323 Κ in water 2.5 16.5 2 Quenching + cooling to 77 Κ (γ → α transformation) + slow heating

(0.2 K/min) to 703 Κ and subsequent rapid heating (200 K/min) from 703 to 873 Κ (α → γ transformation)

7.9 15.6

3 Quenching + slow heating to 743Κ + rapid heating to 873Κ 10.5 15.6 4 Quenching + slow heating to 763Κ + rapid heating to 873Κ 6.0 15.6 5 Quenching + slow heating to 733Κ + rapid heating to 1073Κ 2.4 16.0

Acknowledgments

This work was supported in part by the Russian Foundation for Basic Research (Project no.10-03-00113 – Project no.04-02-96089); and the Presidium of RAS (Project no.12-P-2-1043). References

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