Strength of ultrafine-grained corrosion-resistant steels after severe plastic deformation

Materials Science and Engineering A 387–389 (2004) 244–248

Strength of ultrafine-grained corrosion-resistant steels after severe plastic deformation O.V. Rybal’chenko a,b,∗ , S.V. Dobatkin a,b , L.M. Kaputkina b , G.I. Raab c , N.A. Krasilnikov c a

Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Leninsky Prospect 49, 119991 Moscow, Russia b Moscow State Steel and Alloys Institute, Leninsky Prospect 4, 119049 Moscow, Russia c Institute of Physics of Advanced Materials, Ufa State Aviation Technical University, K. Marx 12, 450000 Ufa, Russia Received 25 August 2003; received in revised form 12 March 2004; accepted 14 March 2004

Abstract It is shown that severe plastic deformation of corrosion resistant Cr–Ni austenitic steels leads to formation of the nano and submicrocrystalline structure. After high-pressure torsion average grain size of about 50 nm is obtained, after equal channel angular pressing—generally oriented structure with separated grains (100–250 nm in size). The severe plastic deformation by both high-pressure torsion and equal channel angular pressing methods promotes the martensitic transformation. Structure of Cr–Ni austenitic steels obtained by severe plastic deformation exhibits a substantial strain hardening. © 2004 Elsevier B.V. All rights reserved. Keywords: Austenitic steel; Severe plastic deformation; Equal channel angular pressing; High-pressure torsion; Nano and submicrocrystalline structure

1. Introduction At present, a great attention is paid to the processes of severe plastic deformation (SPD) due to the opportunity of the formation of nano and submicrocrystalline structures upon deformation [1–3]. Different structures can be obtained, depending on experimental scheme. The limiting structural states are generally realized upon high-pressure torsion (HPT) since, in this case, the applied pressure (up to 10 GPa) allows one to reach a high strain degree [4]. An equal channel angular pressing (ECAP) as one of the most advantageous SPD methods allows to prepare nano and submicrocrystalline samples as large as of 20–40 mm in diameter and 100–150 mm in length [2,5,6]. The pieces of such size can be widely used for medical tools and implants; in particular, they are already tested for titanium [2]. The purposes of the present work were to study the opportunity of the nano and submicrocrystalline structure formation in Cr–Ni austenitic steels upon SPD and to determine the mechanical characteristics of such materials.

∗

Corresponding author. Tel.: +7 095 1357743; fax: 7 095 1358680. E-mail address: [email protected] (O.V. Rybal’chenko).

0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.03.097

The Cr–Ni austenitic steels were chosen for the opportunity of their application in medicine.

2. Experimental procedure The 0.08%C–18.3%Cr–9.8%Ni–0.6%Ti (1) and 0.07%C– 17.3%Cr–9.2%Ni–0.7%Ti (2) austenitic steels were used for the study (Table 1). The first steel was taken after hot rolling, and second one was used after quenching from a temperature of 1050 ◦ C (after holding for 30 min). Deformation of 0.08%C–18.3%Cr–9.8%Ni–0.6%Ti steel was performed by high-pressure torsion on Bridgeman type device [4]. The samples of 10 mm in diameter and 1 mm in thickness had previously been 50% deformed by compression and later by torsion to different strain degrees up to nine revolutions which corresponds to logarithmic strain degree on a half of the sample radius e = 6.4 according to equation [7]. The 0.07%C–17.3%Cr–9.2%Ni–0.7%Ti steel samples of 20 mm in diameter and 80 mm in length were subjected to ECAP by Bc route at room temperature for four passes. The angle between the channels was 90◦ for first pass and

O.V. Rybal’chenko et al. / Materials Science and Engineering A 387–389 (2004) 244–248 Table 1 Chemical composition of steels studied Steels

Mass (%) C

Cr

Ni

Cu

Ti

Si

Mn

S, P

(1) (2)

0.08 0.07

18.3 17.3

9.8 9.2

0.1 0.2

0.6 0.7

0.6 0.6

0.3 1.4

0.003 0.003

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M-400-H “Leco” hardness testing machine with 50 g load. The mechanical properties were determined using an INSTRON 1196 testing machine at a cross-head speed of 1.5 mm/min on the samples of 3 mm in diameter and 18 mm in length of the working area.

3. Results and discussion 120◦

for another three passes. The total true deformation was about e = 3.2 [8]. The samples were annealed on different temperatures up to 900 ◦ C with holding time 30 min. The structure analysis has been carried out using a light microscope “Olympus PME 3”and transmission electron microscope JEM-100CX. The samples were examined by X-ray diffraction using a DRON-1UM diffractometer. The phase composition was determined using a PHAN% program of the quantitative phase analysis taking into account the texture formation. Microhardness was defined using the

3.1. Deformation by HPT As was noted earlier [9], room-temperature deformation by high-pressure torsion (HPT) leads to the formation of separated structure elements with high-angle boundaries already at e = 4.3 (one revolution). As a whole, the oriented structure, which is formed at the initial stages, is transformed into a rather equiaxed structure upon further deformation. Fig. 1a shows such structure for the 0.08%C–18.3%Cr–9.8%Ni–0.6%Ti steel after deformation

Fig. 1. Structure of steel after high pressure torsion: (a) e = 5.8, T = 20 ◦ C; (b) e = 5.8, T = 20 ◦ C; heating at T = 500 ◦ C; (c) e = 5.8, T = 20 ◦ C; heating at T = 700 ◦ C; (d) e = 5.8, T = 20 ◦ C; heating at T = 800 ◦ C.

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by HPT to e = 5.8 (five revolutions). An average size of structural elements is about of 50 nm. The character of the selected electron-diffraction pattern (SAED) generally indicates a high-angle misorientation at the boundaries. Diffracted beams lie randomly around rings demonstrating that this microstructure consists partially of grains separated by boundaries having high angles of misorientation. Therefore, we can define the obtained structure as nanocrystalline. However, we should note, that separated regions of oriented structure were observed even at e = 6.4 (nine revolutions). We revealed that deformation induces a martensitic transformation [9]. The martensite content in the sample was 50% already at e = 4.3 (one revolution) and ∼60% at e = 5.8 (five revolutions). We revealed not only ␥ → ␣ , but also ␥ → ε → ␣ transformation [9]. The X-ray diffraction data on the volume fraction of martensite were obtained without taking into account for texture [9]. As we determined the martensite content with allowance for texture formed upon deformation by torsion [10], the same samples after e = 5.8 (five revolutions) revealed 80% rather than 60% martensite shown earlier [9]. In general, we note that the difference in the martensite contents in austenitic steels subjected to SPD [11,12] is caused not only by the deformation scheme and applied pressure, but also, to the greater extent, by the technique of the determination of the ␣ -phase content. In any case, SPD leads to the formation of a two-phase austenitic–martensitic structure, which should increase the thermal stability of the obtained nanocrystalline steel. The nanocrystalline structure after SPD is characterized by the following specific features: (1) a small grain size, down to nanolevel; (2) mainly high-angle boundaries; (3) a small density of dislocations inside grains and (4) nonequilibrium grain boundaries with a high density of grain-boundary dislocations and stress fields near the boundaries [2]. It is apparent that such structure is unstable upon heating because of a high-density of nonequilibrium grain-boundary surfaces. In this case, the second phase can suppress the matrix phase growth. Upon heating of the nanocrystalline 0.08%C–18.3%Cr– 9.8%Ni–0.6%Ti steel after SPD by HPT, the initial grain size of 50 nm remains virtually unchanged up to a temperature of 400 ◦ C. The grain size slightly increases (to 250 nm) at 500 ◦ C and begins to intensely grow at temperatures above 600 ◦ C (Figs. 1 and 2). This corresponds to the changes in the volume fractions of phase constituents upon heating (Fig. 3). The martensite fraction begins to decrease upon heating above 400 ◦ C. After heating to 550 ◦ C, the phase composition corresponds to a ratio 50:50%. This still suppresses an intense grain growth, which begins at 600 ◦ C, when the austenite content is ∼80%. Upon heating of the nanocrystalline steel to 600 ◦ C, the grain size is retained in a submicrocrystalline range, remaining below 1 ␮m. After heating to 800 ◦ C, the grain size was determined by metallographic examination to be ∼7 ␮m (Figs. 1 and 2).

Fig. 2. Dependence of grain size on the annealing temperature after HPT.

The microhardness evolution upon heating virtually corresponds to the above results. Upon heating to 400 ◦ C, the microhardness remains unchanged (Fig. 4). At higher temperatures, the microhardness decreases due to the grain-boundary recovery, a small grain growth and, probably, a decrease in the fraction of martensite (Figs. 3 and 4). However, as was shown in [9], the difference in the microhardnesses of phase constituents (austenite, ferrite, and martensite) in nanocrystalline corrosion-resistant steels after SPD becomes insignificant because the volume fraction of grain boundaries is very high. 3.2. Deformation by ECA pressing To determine the mechanical characteristics after SPD, we subjected the massive samples to room-temperature de-

Fig. 3. Dependence of phase composition on the annealing temperature after HPT.

O.V. Rybal’chenko et al. / Materials Science and Engineering A 387–389 (2004) 244–248

Fig. 4. Microhardness after SPD and subsequent heating.

formation by ECAP, since the samples deformed by HPT are not suitable for standard mechanical tests. An opportunity to deform a sample in the ECAP die without failure is generally determined by the construction of this die, which is characterized by a decreased friction in the input channel and a back pressure in the output channel. The die used in this work allowed us to deform a sample of 20 mm in diameter and 80 mm in length for four passes i.e. N = 4 (one pass at an angle of 90◦ between channels and three passes at an angle of 120◦ ), to a true deformation e = 3.2. The limiting deformation achieved by ECAP of 0.07%C–17.3%Cr–9.2%Ni–0.7%Ti steel in our work is much lower than that achievable by HPT. For this reason, we failed to obtain an equiaxed structure after ECAP. We observed an oriented structure consisting of elements of 100–250 nm in size (a distance between subgrain or grain boundaries) and some individual equiaxed grains of the same size (Fig. 5). Such oriented structure elements are

Fig. 5. Structure after ECAP N = 4 passes. The plane of the foil is parallel to the axis of the sample.

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presented by shear and deformation bands, twins, martensitic plates, and oriented subgrains (cells) [11]. It is difficult to resolve the structure type in such a fine structure. The oriented structures frequently cross one another at an angle. Sometimes we observed colonies of martensitic plates and/or twins with a spacing of ∼50 nm between boundaries. We observed that the nucleation of equiaxed grains can occur also through the cellular structure. With increasing strain degree, the fraction of the grained structure increases, but even at N = 4 (e = 3.2), the structure remains far from perfection: oriented structural elements, high density of dislocation, large width of boundaries. Unlike HPT, ECAP under the used conditions induces a weak martensitic transformation, which becomes more active only at N = 4, leading to the formation of 45% martensite (Fig. 6). The imperfect structure obtained upon ECAP is characterized by a smaller microhardness (4–4.5 GPa) compared with that of the nanocrystalline structure obtained after HPT (Fig. 4). Upon heating of the steel after ECAP, its microhardness begins to decrease at 500 ◦ C, i.e., at a higher temperature than after HPT. This can be probably caused by the absence of the recovery of nanograin boundaries and the corresponding loss in microhardness. Even the imperfect and oriented submicrocrystalline structure of 0.07%C–17.3%Cr–9.2%Ni–0.7%Ti steel after ECAP provides a good combination of mechanical properties. Already at N = 2, the yield strength is 990 MPa at an elongation of 13%. The further deformation up to N = 4 monotonously increases the yield strength up to 1315 MPa at EL = 11%. This shows that, at N = 4, we are still at an unsteady stage of the structure formation. To obtain a perfect nano or submicrocrystalline structure, one should either increase the degree of deformation, or heat the obtained structure. High degree of the achievable deformation and a high pressure used in [12] resulted in a more perfect grained structure with a grain size of ∼100 nm and, corre-

Fig. 6. Martensitic transformation in austenitic steel during ECAP.

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spondingly, a higher plasticity (EL = 27.5%) at a somewhat higher strength (YS = 1340 MPa).

4. Conclusions (1) Room-temperature SPD of the 0.08%C–18.3%Cr– 9.8%Ni–0.6%Ti austenitic steel by HPT leads to formation of the nanocrystalline structure with a grain size of ∼50 nm. (2) SPD by both HPT and ECAP induces the occurrence of martensitic transformation. Upon HPT, the degree of the transformation (80%) is higher than upon ECAP (45%) due to a higher pressure and a higher achievable strain degree. (3) Nanocrystalline structure in the 0.08%C–18.3%Cr– 9.8%Ni–0.6%Ti austenitic steel obtained by cold HPT is stable during heating up to 400 ◦ C, when microhardness decreases due to recovery of grain boundary. Intense grain growth begins at heating higher than 500 ◦ C simultaneously with increasing of ␥-phase fraction. (4) ECAP of 0.07%C–17.3%Cr–9.2%Ni–0.7%Ti steel under the used conditions leads to the formation of the structure, which is generally oriented with a spacing of 100–250 nm between boundaries, but also contains separated grains of the same size. Such structure exhibits a substantial strain hardening to YS = 1315 MPa relative to the initial state (YS = 250 MPa) and plasticity EL = 11%.

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