Thermal stability of nanocrystalline structure in niobium processed by high pressure torsion at cryogenic temperatures

Materials Science and Engineering A 528 (2011) 1491–1496

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Thermal stability of nanocrystalline structure in niobium processed by high pressure torsion at cryogenic temperatures V.V. Popov ∗ , E.N. Popova, A.V. Stolbovskiy, V.P. Pilyugin Institute of Metal Physics, Urals Division of Russian Academy of Sciences, S. Kovalevskaya Str., 18, Ekaterinburg 620990, Russia

a r t i c l e

i n f o

Article history: Received 26 August 2010 Received in revised form 18 October 2010 Accepted 18 October 2010

Keywords: Nanocrystalline structure Thermal stability Grain refinement Grain growth High-pressure torsion Electron microscopy

a b s t r a c t Evolution of structure of Nb subjected to high-pressure torsion (HPT) in liquid nitrogen and further annealing in the temperature range of 100–600 ◦ C has been studied by transmission electron microscopy (TEM). HPT at the cryogenic temperature of 80 K results in the formation of nanocrystalline structure in Nb, with crystallite sizes of about 75 nm and the record-breaking microhardness of 4800 MPa. The structure obtained is stable at room temperature but possesses relatively low thermal stability, namely, it undergoes recrystallization at lower temperatures than the structure after the conventional deformation or room-temperature HPT. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Nowadays nanostructured materials, that is, the materials with an average size of grains or other structural units not larger than 100 nm attract great attention of the materials science researchers due to their uncommon mechanical behavior and unique structure and properties [1]. Particularly, they possess high strength and superplasticity as well as appreciably higher diffusivities than that of conventional polycrystalline materials [2]. As shown in numerous studies, bulk metal materials with submicrocrystalline structure can be obtained by severe plastic deformation (SPD) accomplished by various techniques [3]. Capabilities of grain refinement in various materials by HPT are demonstrated in a number of recent publications [4–8]. However, it is only very seldom when one obtains a true nanocrystalline structure in pure metals by SPD, as there is a saturation state when further refinement of a material is impossible [9–11]. The smallest grain sizes achievable at HPT depend, first of all, on the type of a material processed, i.e. on its crystallite lattice and melting temperature [9]. Besides, such parameters of HPT as strain, pressure and temperature are of great importance. For example, as shown in [10], the resulting size of structural elements in high purity copper decreases with increasing pressure, and an increasing temperature leads to coarser

∗ Corresponding author. Tel.: +7 343 378 38 41; fax: +7 343 374 52 44. E-mail address: [email protected] (V.V. Popov). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.10.052

structure, the latter effect being much stronger. On the opinion of Pippan et al. [11], the steady-state microstructure in the saturation region requires equilibrium in the generation and annihilation of defects: vacancies, dislocations, high- and low-angle boundaries, and the key process limiting the refinement achievable by SPD is grain boundary migration. As shown in [12,13], single-crystalline high-purity niobium subjected to HPT by 5 revolutions of anvils at room temperature acquires the structure which may be considered as borderline between the nano- and submicrocrystalline with crystallites sizes of 100–120 nm, high dislocation density and wide defective highangle boundaries. Besides, some areas of cellular structure are also retained in it after such treatment. With increasing strain up to 10 revolutions of anvils we did not observe further grain refinement, and it was suggested that at room-temperature HPT the saturation stage is reached in niobium in this strain range, and it undergoes a sort of “dynamic recrystallization”, i.e. its structure restores by grain boundary migration as it occurs in case of dynamic recrystallization during hot working. Thus, at room-temperature HPT the saturation stage is achieved due to the processes similar to dynamic recovery and dynamic recrystallization even in such refractory metal as niobium, and one of the ways to suppress these processes is decreasing deformation temperature. The main goal of the present study is analyzing possibilities of Nb nanostructuring by HPT at cryogenic temperatures (the temperature of liquid nitrogen) and investigation of thermal stability of the structure obtained.

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Fig. 1. Structure of Nb after 1 revolution HPT at 80 K: (a, b) specimen’s center; (c, d) at the radius middle; (e, f) specimen’s edge; (a, c, e) bright-field images with electron diffraction patterns (in (a) zone axis is [1 1 3]); (b, d, f) dark-field images in (1 1 0)Nb reflections.

2. Experimental In the present study high-pressure torsion in Bridgman anvils was carried out in liquid nitrogen. The device construction enables to put the HPT die-set in liquid nitrogen bath and to carry out deformation directly in it. Before the HPT the die-set was held in liquid nitrogen for 30 min. The temperature of a specimen was controlled by a thermo-couple connected to an anvil near the specimen, and it was 78–80 K. Disks of 99.99% pure single-crystalline niobium the diameter of 6 mm and the thickness of 0.5 mm were deformed at the rate of 0.3 revolutions per minute under the load of 8 GPa by 1 and 3 revolutions. The true strain e was calculated as the sum of the true shear strain (esh ) and the true compression strain (ecomp ). The true shear strain was calculated as esh = ln(1 + 2 )1/2 , where  is the shear strain at torsion ( = (ϕR)/h, where ϕ is torsion angle in radians, h is specimen thickness in mm and R is the radial distance from the rotation axis in mm). The true compression strain was calculated as ecomp = ln(h0 /hf ), where h0 and hf are the initial and final specimen thicknesses respectively. The validity of this approach to calculation of the true strain at HPT is well substantiated in [14], though there are different ways of e determination at this technique of SPD [10,11,15]. The as-calculated true strain at the middle of radius is e = 3.9 at one revolution HPT and 7.1 at three revolutions. The as-deformed Nb was annealed in a vacuum furnace at 100, 200, 300, 400, 500 and 600 ◦ C for 1 h to determine thermal stability of the structure obtained. The structure of as-deformed and annealed

specimens was studied by TEM in JEM-200CX and Philips-CM30 electron microscopes with subsequent treatment of images by the computerized program SIAMS-600 [16]. Histograms of grain size distribution were constructed. Microhardness was measured by a special unit in the optical microscope Neophot-21, at the load of 50–100 g, and calculated as H = 18192·P/C2 , MPa, where P is load in g, and C is the indentation diagonal in ␮m. Every value of C was calculated as an average of not less than 9 indentations. 3. Results and discussion As well documented in the literature [see, e.g., 14,17–19], grain refinement at HPT proceeds through three main stages, the first of which is the formation of the dislocation cell structure. With strain growth dislocation density increases, and dislocations are not randomly stored in the structure but concentrate mainly into cell boundaries. At the second stage, with further strain increasing, a mixed structure is formed, containing both cells with somewhat smaller sizes and subgrains with growing misorientation angles between them. And, finally, the third stage is characterized by uniform submicro- or nanocrystalline structure with mainly highangle boundaries between crystallites. After 1 revolution HPT in liquid nitrogen all three stages of structure evolution are observed in Nb (Fig. 1). In the specimen center dislocation cell structure with high dislocation density and markedly non-uniform dislocation distribution is observed. In Fig. 1a, b one can see coarse cells the size

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Fig. 2. Structure of Nb (at the radius middle) after 3 revolutions HPT at 80 K: (a) bright-field image with a ring-wise electron diffraction pattern; (b) dark-field image in (1 1 0)Nb reflection.

of 0.5–0.9 ␮m, with wide dislocation boundaries, and reflections of only one plane are observed in the electron diffraction pattern. At the radius middle mixed structure is observed, and in the electron diffraction pattern several spots located near each other in the Debye rings indicate the formation of low-angle boundaries (Fig. 1c, d). And, finally, on the specimen’s edge after this treatment the nanocrystalline structure is clearly seen, and the electron diffraction patterns are ring-wise, with randomly located bright spots all over the Debye rings (Fig. 1e, f). Such non-uniformity of structure along a specimen radius is characteristic of HPT deformed specimens because of high strain gradient from the center to periphery [20,21]. However, this non-uniformity decreases markedly with growing applied pressure, and, for example, in Ni processed by HPT with the pressure of 9 GPa the structure obtained in specimens center and periphery is practically the same [22]. An analogous result was obtained on commercial pure aluminum in [23]: lower microhardness and less grain refinement were observed in the central regions of the disks in the initial stages of HPT, but the microstructures became reasonably homogeneous across the disks at higher imposed strains. In the case under consideration, i.e. in Nb after 1 revolution of HPT in liquid nitrogen, we observe the formation of crystallites with high-angle boundaries only along the specimen’s edge, where the strain is maximal. However, a specific strained state with high level of internal stresses arises after this treatment throughout the specimen, and this is testified by high values of microhardness, which equals to 2200 MPa in the center and 2900 MPa on the edge of a specimen. With increasing strain at 80 K up to 3 revolutions of anvils (e = 7.1 at the radius middle) the structure changes drastically. It becomes nanocrystalline throughout the specimen’s cross-section, and both the average grain size (75 nm) and the grain size scattering (30–180 nm) are considerably lower than that attained at room temperature HPT [13,24]. It is interesting to note that in this case the structure is quite uniform along the specimen’s radius. A specific curved contrast characteristic of HPT treated materials is observed in grains bulk, and all the electron diffraction patterns are ring-wise, that is, they consist of bright point-wise reflections, not elongated in azimuth directions and located along entire Debye rings (Fig. 2: here and further the images taken at the radius middle are shown). The latter reliably demonstrate that most of crystallite boundaries are high-angle, with random misorientation between grains. The microhardness of this specimen reaches the recordbreaking for Nb value of 4800 MPa, and its value is practically the same along the specimen’s radius contrary to the above-considered specimen after 1 revolution HPT. The value obtained is about twice as high as of niobium nanostructured at room temperature [13] or of high-strength heavily-drawn Cu–Nb composites [25,26] in which anomalously high strength is due to nano-scaled sizes of Nb filaments and interspaces between them in copper matrix [27].

The results obtained are in agreement with Sevillano who analyzed capabilities of SPD for nanostructuring and strengthening and gave a new logical explanation for the more pronounced strengthening at HPT compared to that attained by other techniques of SPD [28]. According to [28], at HPT the effect of high pressure on the structural evolution is summed up with the strain gradient hardening. The latter induces the storage of geometrically necessary dislocations difficult to recover, which results in an extra strengthening and in a more effective nano-domain building than the statistically stored dislocation density promoted by, for example, equal-channel angular extrusion (ECAE). This extra strengthening is akin to that shown by BCC or HCP materials in axisymmetric elongation by wire-drawing [29]. One of the most important problems with submicrocrystalline and nanostructured materials is their lower thermal stability compared to that of ordinary polycrystals which is due to high internal stresses and non-equilibrium boundaries in the former [30–33]. Thermal stability of Nb at 400–800 ◦ C annealing after the room temperature HPT was studied in [12,13], and it was demonstrated that indeed the structure is less thermally stable compared to conventional highly-deformed state, and with increasing strain (from 5 to 10 revolutions of anvils) the thermal stability decreases. The results of studies of thermal stability of the nanocrystalline structure obtained by 3 revolutions HPT in liquid nitrogen are shown in Table 1 and Figs. 3–6. It is important to note that the as-obtained structure appeared to be stable at room temperature, and its 6 months aging did not result in either grain growth, or microhardness drop. Consequently, one can conclude that niobium nanostructured at cryogenic temperature does not undergo recrystallization at ambient temperature contrary to copper in which we also managed to refine the structure up to the nanocrystalline state by HPT in liquid nitrogen, but it appeared to be extremely unstable and undergone recrystallization during a very short period, from several minutes to several days dependently on the strain degree [34]. Actually, that is not surprising taking into account that room temperature is 0.22Tm for copper but only of about 0.1Tm for Nb.

Table 1 Parameters of grain size distribution (D is the grain size, and RMS is the root mean square deviation) and microhardness (H) of Nb after HPT at 80 K by 3 revolutions and annealing. Parameter/treatment

H, MPa

Dav , nm

Dmin , nm

Dmax , nm

RMS, nm

3 rev. HPT at 80 K HPT (6 months later) 3 rev. HPT + 100 ◦ C, 1 h 3 rev. HPT + 200 ◦ C, 1 h 3 rev. HPT + 300 ◦ C, 1 h 3 rev. HPT + 400 ◦ C, 1 h 3 rev. HPT + 500 ◦ C, 1 h 3 rev. HPT + 600 ◦ C, 1 h

4800 4800 4130 3690 3600 3440 1630 1150

75 75 80 90 115 210 >1000 >1000

30 30 35 40 55 80 – –

170 180 130 190 380 470 – –

20 22 17 25 44 65 – –

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100 ◦ C annealing are practically the same as in the as-deformed state, and at the 200 ◦ C annealing they are somewhat larger. Nevertheless, in both cases it is still the nanocrystalline structure, and in most of grains one can see the curved contrast characteristic of the strained state, and all the electron diffraction patterns are ring-wise, with many reflections in the Debye rings (Fig. 5). Thus, the observed microhardness decreasing may result from recovery processes and the decrease of dislocation density. With the annealing temperature increasing up to 300 ◦ C the further grain growth is observed, and along with the dispersed crystallites (the size of less than 100 nm) there appear relatively coarse grains (the size of more than 300 nm) with low dislocation density, the boundaries of which cross at the angle of about 120◦ which is characteristic of recrystallized grains (Fig. 6a). Thus, at as low temperature as 300 ◦ C in some areas of the specimen, the relative fraction of which is not big yet, the grain growth starts, and the grain size scattering in this case gets markedly wider (Fig. 4d). As seen from Fig. 3, the microhardness decreases only slightly at 300 ◦ C annealing because the relative fraction of recrystallized areas is small. The further decrease of microhardness is observed at the 400 ◦ C annealing when further development of recrystallization occurs, namely, the fraction of recrystallized grains increases and the average grain size grows up to 210 nm, i.e. the structure is no more nanocrystalline but it is submicrocrystalline (Figs. 4e and 6b).

Microhardness, MPa

5000

4000

3000

2000

1000 HPT

200

400

600

Temperature,0C Fig. 3. Microhardness of Nb after HPT by 3 revolutions at 80 K and further annealing versus the annealing temperature.

At as low annealing temperature as 100–200 ◦ C the microhardness decrease starts (Fig. 3) whereas the structure and grain sizes change only slightly (Table 1 and Figs. 4 and 5). As seen from Fig. 4a–c, the mean grain size and grain size scattering after the

100

The number of grains

a

80 60 40 20 0 0

40

80 120 D, nm

160

80 b

80 60 40 20 0 0

200

40

80

160

200

40 20 0 0

40

80

120

160

200

D, nm

30

The number of grains

d

80 60 40 20 0 0

120

c

60

D, nm

100

The number of grains

The number of grains

100

The number of grains

1494

100

200 300 D, nm

400

500

e 20

10

0 0

100

200

300

400

500

D, nm

Fig. 4. Histograms of grain size distribution in Nb after HPT by 3 revolutions at 80 K (a) and annealing at 100 ◦ C (b), 200 ◦ C (c), 300 ◦ C (d), and 400 ◦ C (e).

Fig. 5. Structure of Nb (at the radius middle) after 3 revolutions HPT at 80 K and annealing: (a) 100 ◦ C, 1 h; (b) 200 ◦ C, 1 h.

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Fig. 6. Structure of Nb (at the radius middle) after 3 revolutions HPT at 80 K and annealing: (a) 300 ◦ C, 1 h; (b) 400 ◦ C, 1 h; (c) 500 ◦ C, 1 h, zone axis [3 3 1]; (d) 600 ◦ C, 1 h, zone axis [1 1 0].

It should be noted that in this case the ring-wise electron diffraction pattern transforms into the point-wise, i.e. there are no many randomly distributed spots in the Debye rings in this case, but only 2–3 spots (compare inserts in Figs. 5 and 6). At higher annealing temperature, beginning from 500 ◦ C, the intensive grain growth is observed. Crystallite sizes reach 1 ␮m or even more, and on the electron diffraction patterns one can see reflections of only one plane of the reciprocal lattice which means that the structure becomes microcrystalline (Fig. 6c, d). In this case microhardness drastically drops to the same values as were observed after the room-temperature HPT by 10 revolutions and annealing at 600–700 ◦ C [13]. Thus, it may be stated that beginning from 500 ◦ C annealing the structure recrystallizes completely throughout the whole specimen, the crystallites sizes growing with the growth of the annealing temperature. Nevertheless, the values of microhardness are still higher than that typical of niobium after cold drawing and further annealing (800 MPa). An analogous result was obtained, for example, in the studies of thermal stability of chromium after SPD [35]. As found in that paper, nanostructured Cr at 500 ◦ C annealing undergoes non-uniform recrystallization with drastic microhardness drop. At 700 ◦ C uniform grains coarsening is observed in Cr with further microhardness drop, but, nevertheless, the latter is higher compared to that in the specimens which were not subjected to SPD. These data demonstrate that at SPD by high-pressure torsion not only grain refinement to nano-scaled sizes occurs, but a specific stressed state arises which makes the material different from an ordinary polycrystal. It is of interest to compare thermal stability of structure after HPT at room temperature and in liquid nitrogen. As shown in [12,13], the structure of specimens subjected to room temperature HPT by 5 revolutions of anvils is stable up to 600 ◦ C. With the strain degree growth up to 10 revolutions further grain refinement is not observed, but the structure gets less stable, and the intensive grain growth starts at 500 ◦ C. As shown in the present study, in the specimens deformed in liquid nitrogen by 3 revolutions the intensive

grain growth is observed at as low temperatures as 400 ◦ C. Thus, the decrease of HPT temperature, on the one hand, allows to obtain the true nanocrystalline structure (that is, the achievable grain refinement is more pronounced at cryogenic temperature than at room temperature HPT and the mean grain size is smaller in the former case). This observation is in agreement with [11], where it is demonstrated on a number of various materials that the saturation grain sizes strongly depend on HPT deformation temperature and decrease with temperature drop, though it is noted that this effect is less pronounced at low temperatures than at medium temperatures. On the other hand, the nanocrystalline structure obtained in Nb by HPT in liquid nitrogen is less thermally stable and its degradation starts at as low as 300 ◦ C, though it is quite stable at room temperature ageing.

4. Conclusion TEM investigations carried out in the present study have shown that at cryogenic temperatures HPT results in the formation of true nanocrystalline structure in niobium, namely, the structure consists of grains with high-angle boundaries, the average size of 75 nm, with narrow grain size scattering. This structure possesses high level of internal stresses and high dislocation density which is evidenced by the specific contrast in electron micrographs as well as by the record-breaking value of microhardness (4800 MPa). The as-obtained structure is quite stable at room temperature, i.e. no grain growth is observed on heating up to the ambient temperature and further holding at this temperature for several months. However, at annealing the thermal stability of this structure appears to be lower compared to the structure obtained by room temperature HPT. In nanocrystalline Nb obtained by HPT in liquid nitrogen, some grain boundaries migrate at as low temperatures as 300 ◦ C, at 400 ◦ C a marked grain growth occurs, and the structure turns into submicrocrystalline, and complete recrystallization is observed at 500 ◦ C. The low thermal stability of the structure in Nb subjected

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to HPT at cryogenic temperature is obviously due to high density of defects and to the pronounced non-equilibrium state of grain boundaries. Acknowledgements The work has been done on the equipment of Collective Use Electron Microscope Center of IMP Urals Division of RAS within the Presidium of RAS program with partial support of Russian Foundation on Basic Research (project No. 10-03-00530). References [1] H. Gleiter, Acta Mater. 48 (2000) 1–29. [2] R.Z. Valiev, T.C. Lowe, A.K. Mukherjee, JOM: J. Miner. Met. Mater. Sci. 52 (4) (2000) 37–40. [3] R.Z. Valiev, I.V. Aleksandrov, Bulk Nanostructured Metallic Materials: Manufacturing, Structure and Properties, ACADEMKNIGA, Moscow, 2007, 398 pp (in Russian). [4] R.K. Guduru, K.L. Murty, K.M. Youssef, R.O. Scattergood, C.C. Koch, Mater. Sci. Eng. A463 (1–2) (2007) 14–21. [5] K. Edalati, Z. Horita, Y. Mine, Mater. Sci. Eng. A527 (7–8) (2010) 2136–2141. [6] Y. Ito, Z. Horita, Mater. Sci. Eng. A503 (1–2) (2009) 32–36. [7] N. Lugo, N. Llorca, J.M. Cabrera, Z. Horita, Mater. Sci. Eng. A477 (1–2) (2008) 366–371. [8] A.P. Zhilyaev, A.A. Gimazov, E.P. Soshnikova, Á. Révész, T.G. Langdon, Mater. Sci. Eng. A489 (1–2) (2008) 207–212. [9] A.V. Korznikov, A.N. Tyumentsev, I.A. Ditenberg, Phys. Met. Metallogr. 106 (4) (2008) 418–423. [10] T. Hebesberger, A. Vorhauer, H.P. Stuwe, R. Pippan, Proc. Conf. “Nanomaterials by Severe Plastic Deformation-NANOSPD2”, Vienna, Austria, 2002, pp. 447–452. [11] R. Pippan, S. Scheriau, A. Taylor, et al., Annu. Rev. Mater. Res. 40 (2010) 319–343. [12] E.N. Popova, V.V. Popov, E.P. Romanov, V.P. Pilyugin, Phys. Met. Metallogr. 101 (1) (2006) 52–57.

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