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Microstructure evolution and mechanical properties of nanocrystalline zirconium processed by surface circulation rolling treatment
Materials Science & Engineering A 565 (2013) 27–32
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure evolution and mechanical properties of nanocrystalline zirconium processed by surface circulation rolling treatment Chao Yuan a,b, Ruidong Fu a,b,n, Fucheng Zhang a,b, Xiangyi Zhang a,b, Fengchao Liu c a
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, Hebei 066004, PR China College of Materials Science and Engineering, Yanshan University, Qinhuangdao, Hebei 066004, PR China c School of Mechanical & Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore b
a r t i c l e i n f o
abstract
Article history: Received 11 October 2012 Received in revised form 13 November 2012 Accepted 15 November 2012 Available online 12 December 2012
In this work, the microstructure evolution and nanoindentation hardness of nanocrystalline zirconium processed by surface circulation rolling treatment at cryogenic temperatures were investigated in details. Experimental results indicated that the total deformation layer depth exceeds 600 mm. The average grain sizes vary from about 8 nm in the topmost surface to micrometers in the coarse grained matrix, corresponding to a gradient variation in hardness from about 6.0 to 2.86 GPa. The microstructure evolutions were found that the deformation bands form at the initial stage of the deformation. With increasing the strain, the dislocation cells form in interior of the deformation bands and finally transformed into nanograins. The Hall–Petch relationship between hardness and grain size is not linear due to the change of the deformation mechanism from dislocation pile-ups to grainboundary sliding as the grain size becomes smaller. & 2012 Elsevier B.V. All rights reserved.
Keywords: Nanocrystalline materials Microstructure Nanoindentation hardness Surface circulation rolling treatment (SCRT) Zirconium
1. Introduction Nanocrystalline materials exhibit many excellent properties, such as excellent hardness and strength, outstanding tribological and magnetic properties, and improved toughness, compared with conventional polycrystalline materials [1–5]. Consequently, these materials have been extensively researched in the past decades. However, difficulties in preparing bulk nanocrystalline materials limit their engineering applications. For metals, most failures originate from the surface, whose surface microstructures and properties are very sensitive. Hence, global performances of metallic materials can be improved by surface modification [6]. Lu and Lu [6] first proposed the concept of surface nanocrystallization (SNC) of metallic materials. Since then, fabrication of nanostructured layer on a bulk material surface has received increasing attention. Various mechanical processing techniques have been successfully used to fabricate nanostructures on metallic material surfaces. Such techniques include surface mechanical attrition treatment (SMAT) [7,8], high-energy shot peening [9], ultrasonic shot peening [10], surface mechanical grinding treatment (SMGT) [11,12], and so on. Meanwhile, the performance of nanostructured layers has also been investigated. n Corresponding author at: State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, Hebei 066004, PR China. Tel.: þ86 335 807 4792; fax: þ 86 335 807 4545. E-mail address: [email protected] (R. Fu).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.11.092
For example, an ideal gradient nano-micro-grained architecture was synthesized in a bulk coarse-grained Cu substrate. The yield strength of the gradient nano-/coarse-grained structure was twice that of the coarse-grained substrate. The uniform elongation between these substrates was exactly the same [12]. These results were desirable and profoundly influenced subsequent studies. The grain refinement mechanism in the SNC process is similar to that in severe plastic deformation (SPD) techniques, extensively used to refine grains in various metals and alloys over the past decades [13,14]. The mechanism of SPD-induced grain refinement can usually be attributed to phase transformation, dislocation activities, and/or deformation twinning. Generally, crystal structure, stacking fault energy (SFE), and initial grain size were considered as the intrinsic factors that affect the above processes [15]. For example, for bcc Fe with a high SFE, the formation of dislocation walls and dislocation tangles resulted in the grain refinement [7]; for fcc Cu with a medium SFE, grains were refined via formation of dislocation cells, dislocation walls and twins [16]; while for hcp Ti with a high SFE and Mg with a medium SFE, the grain refinement process involved the formation of twins, dislocation arrays and dense dislocation walls (DDWs) [17,18]. On the other hand, large compression and/or shear strain, high strain rates, and low temperature were widely considered as the extrinsic factors that govern the grain refinement process. For example, the two most important SPD techniques, namely equalchannel angular pressing and high-pressure torsion, mainly draw large shear strain from the deformation process at higher
C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
3.1. X-ray diffraction analysis Fig. 2 shows the XRD patterns of the surface layer of the SCRTed and un-SCRTed samples. The Miller indices of the reflecting planes corresponding to each peak are given in the figure. It is noted that SCRT does not result in any phase transformation of Zr; however, evident Bragg-diffractional peak broadening exists in the SCRTed Zr, which is caused by the grain refinement and an increase in the atomic-level microstrain [22]. The average grain size calculated by the Scherrer and Wilson method [23] was about 42 nm, and the microstrain was about 0.18%.
(002)
un-SCRT SCRT
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65
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75
Fig. 2. XRD patterns of the SCRTed and un-SCRTed samples.
ω
υ
r Deformed layer Sample
Fig. 1. Schematic illustrations of (a) the tool tip outline and (b) the SCRT process.
(202)
(004)
(201)
(103)
50
2-theta (degree)
Tool tip
d
(200) (112)
(110)
(102)
A commercial pure zirconium (Zr 702) plate (90 mm 65 mm 3 mm) was used for the SCRT experiment. The plate was annealed at 1223 K for 4 h to obtain a coarse-grained polycrystalline structure. The average grain size of the as-annealed sample is about 60 mm. Fig. 1 shows a schematic illustration of the SCRT set-up used in the present work. A cylindrical tool with a curved surface at the tip (Fig. 1a) rotates at rotational speed o and maintains a penetration depth d into the sample surface. The sample (or tool) moves at a travel speed u along the horizontal direction. The plastic strain and depth of the deformation layer depend on the penetration depth and the tool tip curvature radius r (Fig. 1b). In order to refine the original coarse grains into nanograins, effective cooling of the processed material and the tool is necessary during SCRT. Before the SCRT the plate sample and the tool were immersed into liquid nitrogen for 5 min in order to ensure effective pre-cooling. During the SCRT the plate sample was soaked in liquid nitrogen all the time. The SCRT processing parameters are described as follows: o ¼400 rpm, u¼ 100 mm min 1, d ¼0.1 mm, r ¼40 mm. In order to effectively refine the grains in the surface layer and increase the thickness of the nanostructured layer, the SCRT process was repeated three times with the same processing parameters. A total reduction of 0.2 mm and a strain rate of about 104–105 s 1 were achieved in the SCRT Zr.
3. Results and discussion
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2. Experimental
The cross-sectional observations of the treated Zr sample were preformed on a HITACHI S-4800 scanning electron microscope (SEM). The microstructure features in the surface layer were characterized by using a JEOL-2010 transmission electron microscope (TEM) at an operating voltage of 200 kV. The cross-sectional hardness and elastic modulus profile of the SCRT surface layer were measured by using a TI900 Triboindenter.
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deformation temperature or lower strain rates [19,20]. However, the formation of nanograins in existing SNC techniques mainly originates from the compression strain at high strain rates and/or cryogenic temperatures. Accordingly, if large shear strain at high strain rates and low temperature are introduced into the SNC process, both the processing method and the grain refining mechanism will undoubtedly be interesting issues. Zirconium (Zr) with an hcp structure has important applications in nuclear industry, aviation and surgical implant due to its resistance to corrosion and irradiation, high melting point and good biocompatibility [21]. In the present work, we chose Zr as a model metal and developed a novel and effective technique, namely surface circulation rolling treatment (SCRT), to fabricate a large-area nanostructured surface on Zr plate at cryogenic temperatures. The microstructure evolutions and hardness of nanocrystalline Zr are investigated in details.
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C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
II
III
Treated surface
I
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Fig. 3. Cross-sectional SEM image of the zirconium sample after SCRT.
3.2. Scanning electron microscopy observations The cross-sectional SEM image of the SCRT sample is shown in Fig. 3. It is obvious that after SCRT, three distinct regions can be recognized: (I) the top surface region (distanced the surface to about 300 mm) which experienced the severe plastic deformation. In this region, the original grain boundaries could hardly be identified under the SEM, indicating surface grain refinement induced by SCRT; (II) the transition region (distanced the surface from about 300 to 600 mm) which was affected by smaller plastic deformation. Comparing to the top surface region, part of the grain boundaries were clearly visible and the grains were elongated to a certain extent; (III) the strain-free coarse-grained region (distanced the surface to about 600 mm) which was not affected during SCRT process. Accordingly, the depth of plastic deformation layer can be deduced to exceed 600 mm, which is higher than that obtained by means of SMAT [7,8] and SMGT [11]. To take insight into the microstructure evolutions of the deformed layers, based on the SEM image, the TEM observations of the microstructures at 400, 250, 100 and 30 mm depths from the surface were carried out. 3.3. Quantitative analysis on microstructure evolution As shown in Fig. 4(a), the microstructure of the layer at the depth of approximately 400 mm from the surface displays parallel deformation bands, which are several hundreds of nanometers in width and a few micrometers in length. These deformation bands divided the Zr coarse grains into finer plates. Meanwhile, many slip systems located in the deformation bands were activated. High-density dislocations and interactions occurred to form dislocation cells and pile up at the boundaries because the boundaries acted as obstacles for dislocation slips [24]. Deduction can be made that deformation bands dominate the plastic deformation at the low strain. It is also consistent with previous report [8,25]. Fig. 4(b) shows the microstructure of the layer at the depth of about 250 mm. Since the relative atomic movement is limited in the deformation bands, the gross deformation processed by deformation bands is quite small, and most of the plastic flow is due to movement of dislocations [26,27]. Hence, with further increase in the strain, the dislocation slips and interactions became the dominant deformation mode instead of deformation bands. The dislocation rearrangement led to the formation of dislocation cells that minimized the total system energy. The similar results are also observed in other hcp metals, such as Ti [17] and Mg [18]. In addition, the intersection between the deformation bands reported in the deformed layers of SMAT [8] was not evident. Compared with SMAT and SMGT, in which only a local small-area shear deformation occurred, the SCRT process produced large-area shear deformation. This deformation mode may be more beneficial to the activation of dislocation slips. As a result, a large number of dislocations emerged within the deformation bands, which further subdivided the plates into small blocks. With increasing strain, the accumulation and
29
rearrangement of dislocations resulted in direct transformation of small blocks into the dislocation cells without the intersection of the deformation bands. The microstructure of the layer at the depth of about 100 mm is shown in Fig. 4(c). Dislocation cells continuously transformed into submicronic grains with the increase in the applied strain. The grain size decreased from more than 200 nm (marked A) to less than 100 nm (marked B). It was also found that dislocations (marked with arrows) pile up inside the grains. Fig. 4(d) shows the microstructure of the layer at the depth of about 30 mm. The grain sizes were less than 100 nm. DDWs (marked with arrows) occurred because of dislocation accumulation and rearrangement. The corresponding SAED pattern with elongated and discrete diffraction arcs indicated that the grain orientations were random, and the submicronic grains shown in Fig. 4(c) started to transform into nanograins. It is believed that the increase of strain is responsible for the subdivision of the subgrains into nanograins [8]. Fig. 4(e) shows a TEM bright-field image taken from the topmost surface of the specimen after SCRT. Under very high strain, the DDWs transformed into grain boundaries, and the submicronic grains were subdivided into equiaxial nanograins. The inset of Fig. 4(e) shows the corresponding SAED pattern with a series of rings, which indicates the random crystallographic orientations of the nanograins. The histogram of the grain-size distribution obtained from a large number of dark-field TEM images was characterized by normal logarithmic distribution with a narrow-size distribution ranging from 4 nm to 30 nm (as shown in Fig. 4f). The average grain size was approximately 8 nm, significantly smaller than that found by X-ray diffraction (XRD) analysis (42 nm). The discrepancy can be ascribed to the fact that the XRD structure information averaged approximately 5 mm thick at the top surface layer [7], as well as to surface roughness. Actually, when the grain sizes are less than a critical value, the dislocation multiplication rate is balanced by the annihilation rate [7]. As a result, dislocations will not accumulate in the grains, and the nanograins become free of dislocations, i.e., the increase in strains could not significantly reduce the nanograin size [7,8]. With increasing depth from the treated surface, grain sizes gradually increase. The distribution of the average structure size determined from a large number of TEM images with the corresponding distance from the treated surface is shown in Fig. 5. An increasing trend is obvious for the grain sizes with increasing depth. A gradient structure is formed with a continuously increasing grain size from nano-scale to the submicron and micron scales in the treated surface layer. Within the top 80 mm thick layer, the grain sizes are less than 100 nm. Hence, a nanostructured layer about 80 mm thick is achieved in the sample by using SCRT. 3.4. Nanoindentation hardness measurements The distributions of hardness and elastic modulus across the treated surface layer are shown in Fig. 6(a). The hardness increased from approximately 2.86 GPa (or 291.84 in the Hv scale) in the coarse-grained matrix to approximately 6.0 GPa at the topmost surface, in accordance with the classic Hall–Petch (H–P) relationship and consistent with the results in references [28] and [29]. The hardness of the top surface layer is about 2 times higher than that of the coarse-grained matrix. We believe that both the intrinsic high hardness and the nano-grains of the surface layer contribute to the observed high hardness of the surface layer of the SCRT specimen [30]. However, this hardness variation trend is much slower than that in previous reports [11,30], which could be more conducive to the performance of the material surface. On the other hand, the elastic modulus maintained a constant
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C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
Fig. 4. TEM micrographs of the layer at the depth of about (a) 400 mm, (b) 250 mm, (c) 100 mm (inset: corresponding SAED pattern) and (d) 30 mm (inset: corresponding SAED pattern) and the top surface layer: (e) bright-field image (inset: corresponding SAED pattern) and (f) dark-field image (inset: histogram of grain size distribution derived from a total of 963 grains).
value of approximately 53.63 GPa over the entire deep range, indicating that the elastic modulus weakly depends on the grain size, also consistent with previous reports [11,31]. Fig. 6b presents the H–P relationship between the nanoindentation hardness, H, and the average grain size, d, revealing that the H–d 1/2 relationship is not linear, but is concave towards to the d 1/2
axis. The hardness as a function of the inverse of the square root of the average grain size is linear at locations 20 mm away from the treated surface. However, within a range of 20 mm from the treated surface, the nanohardness data deviate from the straight line and bend towards the d 1/2 axis. The non-linear H–d 1/2 relationship may due to the change in the deformation mechanism as the average
Average structure size (nm)
C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
grain boundaries. It is consistent with the TEM results above. For small grain sizes (i.e. do17 nm), the dislocations existing in the grains are rare, or even only one. So the H–P relationship is not applicable. Furthermore, the grain boundary regions of nanocrystalline materials will appear relaxation process under the stress, the hardness of materials decreases. The hardness has been dictated by both dislocation interactions and grain-boundary sliding acting simultaneously [32–35]. The transition of the deformation mechanism occurs in the intermediate size of the grain sizes ranging from 17 nm to 25 nm.
100000 10000 1000 250 200 150 100
4. Conclusions
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Distance from surface (µm) Fig. 5. Variations of the average structure size with distance from treated surface.
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A novel and efficient SNC technique, namely SCRT, has been developed to fabricate a nanostructured surface layer on commercial pure zirconium. The thickness of the deformation layer exceeded 600 mm. The average grain sizes varied from about 8 nm in the topmost surface to micrometers in the coarse grained matrix. Meanwhile, a nanostructured layer about 80 mm thick was achieved in the sample. The occurrence of deformation bands was responsible for the grain refinement during the initial stage of deformation. With increasing strain, dislocation cells were formed in the deformation bands and finally transformed into nanograins. The SCRT layer hardness exhibited a gradient variation from approximately 6.0 GPa at the top surface to 2.86 GPa in the coarse-grained matrix, whereas the elastic modulus retained a constant value of approximately 53.63 GPa throughout the entire deep range. However, the Hall–Petch relationship between hardness and average grain size was not linear due to the change of the deformation mechanism from dislocation pile-ups to grainboundary sliding as the grain size became smaller. The SCRT technique can easily manufacture large-area nanostructured layer on the metal plate by a simple process. The improved surface performance is worth to expect.
Distance from surface (μm) Acknowledgments
Distance from surface (μm) 600 400 300 10050
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This work was supported in part by the National Basic Research Program of China (Grant no. 2010CB731606) and in part by the National Science Foundation for Distinguished Young Scholars (Grant no. 50925522).
6.5 60 2 280 120 76 6.0
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References 5.0
[1] [2] [3] [4]
4.5 4.0
[5] 3.5
[6] [7]
3.0 2.5 0
2
4
6
8
10
12
14
Inverse of the square root of the average grain size (μm-1/2) Fig. 6. (a) Variations of hardness and elastic modulus with distance from treated surface and (b) the hardness, H, as a function of the inverse of the square root of the average grain size, d 1/2 Fig. 2.
grain size becomes smaller. For large grain sizes (i.e. d & 25 nm), the hardness is mainly controlled by the increase in the internal stress due to the dislocation pile-ups which results from the presence of
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556. H. Gleiter, Acta Mater. 48 (2000) 1–29. W.L. Li, N.R. Tao, Z. Han, K. Lu, Wear 274–275 (2012) 306–312. K.S. Kumar, H. Van Swygenhoven, S. Suresh, Acta Mater. 51 (2003) 5743–5774. M. Dao, L. Lu, R.J. Asaro, J.T.M. De Hosson, E. Ma, Acta Mater. 55 (2007) 4041–4065. K. Lu, J. Lu, J. Mater. Sci. Technol. 15 (1999) 193–197. N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, Acta Mater. 50 (2002) 4603–4616. L. Zhang, Y. Han, J. Lu, Nanotechnology 19 (2008) 165706. G. Liu, S.C. Wang, X.F. Lou, J. Lu, K. Lu, Scr. Mater. 44 (2001) 1791–1795. G. Liu, J. Lu, K. Lu, Mater. Sci. Eng. A 286 (2000) 91–95. W.L. Li, N.R. Tao, K. Lu, Scr. Mater. 59 (2008) 546–549. T.H. Fang, W.L. Li, N.R. Tao, K. Lu, Science 331 (2011) 1587–1590. R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Prog. Mater. Sci. 45 (2000) 103–189. K. Lu, J. Lu, Mater. Sci. Eng. A 375–377 (2004) 38–45. Y.B. Wang, M. Louie, Y. Cao, X.Z. Liao, H.J. Li, S.P. Ringer, Y.T. Zhu, Scr. Mater. 62 (2010) 214–217. K. Wang, N.R. Tao, G. Liu, J. Lu, K. Lu, Acta Mater. 54 (2006) 5281–5291. K.Y. Zhu, A. Vassel, F. Brisset, K. Lu, J. Lu, Acta Mater. 52 (2004) 4101–4110. H.Q. Sun, Y.N. Shi, M.X. Zhang, K. Lu, Acta Mater. 55 (2007) 975–982. Q. Wei, Z.L. Pan, X.L. Wu, B.E. Schuster, L.J. Kecskes, R.Z. Valiev, Acta Mater. 59 (2011) 2423–2436.
32
C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
[20] G.Q. Zhao, S.B. Xu, Y.G. Luan, Y.J. Guan, N. Lun, X.F. Ren, Mater. Sci. Eng. A 437 (2006) 281–292. [21] D.F. Guo, M. Li, Y.D. Shi, Z.B. Zhang, H.T. Zhang, X.M. Liu, B.N. Wei, X.Y. Zhang, Mater. Des 34 (2012) 275–278. [22] M. Wen, G. Liu, J.F. Gu, W.M. Guan, J. Lu., Surf. Coat. Technol 202 (2008) 4728–4733. [23] H.P. Klug, L.E. Alexander, X-Ray Diffraction Procedures for Polycrystalline and Amorphous Materials, Wiley, New York, 1974, p. 661. [24] L. Lu, Y.F. Shen, X.H. Chen, L.H. Qian, K. Lu, Science 304 (2004) 422–426. [25] W.S. Choi, H.S. Ryoo, S.K. Hwang, M.H. Kim, Metall. Mater. Trans. A 33 (2002) 973–980. [26] T.H. Courtney, Mechanical Behavior of Materials, McGraw-Hill, New York, 1990, p. 309.
[27] R.E. Reed-Hill, Physical Metallurgy Principles, Affiliated East, West Press, New Delhi, 1973. [28] E.O. Hall, Proc. Phys. Soc. London Sect B 64 (1951) 747–753. [29] N.J. Petch, J. Iron Steel Inst. London 174 (1953) 25. [30] P. Jiang, Q. Wei, Y.S. Hong, J. Lu, X.L. Wu, Surf. Coat. Technol 202 (2007) 583–589. [31] T.D. Shen, C.C. Koch, T.Y. Tsui, G.M. Pharr, J. Mater. Res. 10 (1995) 2892–2896. [32] L.L. Shaw, A.L. Ortiz, J.C. Villegas, Scr. Mater. 58 (2008) 951–954. [33] K.S. Kumar, H. Van Swygenhoven, S. Suresh, Acta Mater. 51 (2003) 5743–5774. [34] K.A. Padmanabhan, G.P. Dinda, H. Hahn, H. Gleiter, Mater. Sci. Eng. A (2007) 462–468452-3 (2007) 462–468. [35] C.S. Pande, K.P. Cooper, Prog. Mater. Sci. 54 (2009) 689–706.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure evolution and mechanical properties of nanocrystalline zirconium processed by surface circulation rolling treatment Chao Yuan a,b, Ruidong Fu a,b,n, Fucheng Zhang a,b, Xiangyi Zhang a,b, Fengchao Liu c a
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, Hebei 066004, PR China College of Materials Science and Engineering, Yanshan University, Qinhuangdao, Hebei 066004, PR China c School of Mechanical & Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore b
a r t i c l e i n f o
abstract
Article history: Received 11 October 2012 Received in revised form 13 November 2012 Accepted 15 November 2012 Available online 12 December 2012
In this work, the microstructure evolution and nanoindentation hardness of nanocrystalline zirconium processed by surface circulation rolling treatment at cryogenic temperatures were investigated in details. Experimental results indicated that the total deformation layer depth exceeds 600 mm. The average grain sizes vary from about 8 nm in the topmost surface to micrometers in the coarse grained matrix, corresponding to a gradient variation in hardness from about 6.0 to 2.86 GPa. The microstructure evolutions were found that the deformation bands form at the initial stage of the deformation. With increasing the strain, the dislocation cells form in interior of the deformation bands and finally transformed into nanograins. The Hall–Petch relationship between hardness and grain size is not linear due to the change of the deformation mechanism from dislocation pile-ups to grainboundary sliding as the grain size becomes smaller. & 2012 Elsevier B.V. All rights reserved.
Keywords: Nanocrystalline materials Microstructure Nanoindentation hardness Surface circulation rolling treatment (SCRT) Zirconium
1. Introduction Nanocrystalline materials exhibit many excellent properties, such as excellent hardness and strength, outstanding tribological and magnetic properties, and improved toughness, compared with conventional polycrystalline materials [1–5]. Consequently, these materials have been extensively researched in the past decades. However, difficulties in preparing bulk nanocrystalline materials limit their engineering applications. For metals, most failures originate from the surface, whose surface microstructures and properties are very sensitive. Hence, global performances of metallic materials can be improved by surface modification [6]. Lu and Lu [6] first proposed the concept of surface nanocrystallization (SNC) of metallic materials. Since then, fabrication of nanostructured layer on a bulk material surface has received increasing attention. Various mechanical processing techniques have been successfully used to fabricate nanostructures on metallic material surfaces. Such techniques include surface mechanical attrition treatment (SMAT) [7,8], high-energy shot peening [9], ultrasonic shot peening [10], surface mechanical grinding treatment (SMGT) [11,12], and so on. Meanwhile, the performance of nanostructured layers has also been investigated. n Corresponding author at: State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, Hebei 066004, PR China. Tel.: þ86 335 807 4792; fax: þ 86 335 807 4545. E-mail address: [email protected] (R. Fu).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.11.092
For example, an ideal gradient nano-micro-grained architecture was synthesized in a bulk coarse-grained Cu substrate. The yield strength of the gradient nano-/coarse-grained structure was twice that of the coarse-grained substrate. The uniform elongation between these substrates was exactly the same [12]. These results were desirable and profoundly influenced subsequent studies. The grain refinement mechanism in the SNC process is similar to that in severe plastic deformation (SPD) techniques, extensively used to refine grains in various metals and alloys over the past decades [13,14]. The mechanism of SPD-induced grain refinement can usually be attributed to phase transformation, dislocation activities, and/or deformation twinning. Generally, crystal structure, stacking fault energy (SFE), and initial grain size were considered as the intrinsic factors that affect the above processes [15]. For example, for bcc Fe with a high SFE, the formation of dislocation walls and dislocation tangles resulted in the grain refinement [7]; for fcc Cu with a medium SFE, grains were refined via formation of dislocation cells, dislocation walls and twins [16]; while for hcp Ti with a high SFE and Mg with a medium SFE, the grain refinement process involved the formation of twins, dislocation arrays and dense dislocation walls (DDWs) [17,18]. On the other hand, large compression and/or shear strain, high strain rates, and low temperature were widely considered as the extrinsic factors that govern the grain refinement process. For example, the two most important SPD techniques, namely equalchannel angular pressing and high-pressure torsion, mainly draw large shear strain from the deformation process at higher
C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
3.1. X-ray diffraction analysis Fig. 2 shows the XRD patterns of the surface layer of the SCRTed and un-SCRTed samples. The Miller indices of the reflecting planes corresponding to each peak are given in the figure. It is noted that SCRT does not result in any phase transformation of Zr; however, evident Bragg-diffractional peak broadening exists in the SCRTed Zr, which is caused by the grain refinement and an increase in the atomic-level microstrain [22]. The average grain size calculated by the Scherrer and Wilson method [23] was about 42 nm, and the microstrain was about 0.18%.
(002)
un-SCRT SCRT
30
35
40
45
55
60
65
70
75
Fig. 2. XRD patterns of the SCRTed and un-SCRTed samples.
ω
υ
r Deformed layer Sample
Fig. 1. Schematic illustrations of (a) the tool tip outline and (b) the SCRT process.
(202)
(004)
(201)
(103)
50
2-theta (degree)
Tool tip
d
(200) (112)
(110)
(102)
A commercial pure zirconium (Zr 702) plate (90 mm 65 mm 3 mm) was used for the SCRT experiment. The plate was annealed at 1223 K for 4 h to obtain a coarse-grained polycrystalline structure. The average grain size of the as-annealed sample is about 60 mm. Fig. 1 shows a schematic illustration of the SCRT set-up used in the present work. A cylindrical tool with a curved surface at the tip (Fig. 1a) rotates at rotational speed o and maintains a penetration depth d into the sample surface. The sample (or tool) moves at a travel speed u along the horizontal direction. The plastic strain and depth of the deformation layer depend on the penetration depth and the tool tip curvature radius r (Fig. 1b). In order to refine the original coarse grains into nanograins, effective cooling of the processed material and the tool is necessary during SCRT. Before the SCRT the plate sample and the tool were immersed into liquid nitrogen for 5 min in order to ensure effective pre-cooling. During the SCRT the plate sample was soaked in liquid nitrogen all the time. The SCRT processing parameters are described as follows: o ¼400 rpm, u¼ 100 mm min 1, d ¼0.1 mm, r ¼40 mm. In order to effectively refine the grains in the surface layer and increase the thickness of the nanostructured layer, the SCRT process was repeated three times with the same processing parameters. A total reduction of 0.2 mm and a strain rate of about 104–105 s 1 were achieved in the SCRT Zr.
3. Results and discussion
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2. Experimental
The cross-sectional observations of the treated Zr sample were preformed on a HITACHI S-4800 scanning electron microscope (SEM). The microstructure features in the surface layer were characterized by using a JEOL-2010 transmission electron microscope (TEM) at an operating voltage of 200 kV. The cross-sectional hardness and elastic modulus profile of the SCRT surface layer were measured by using a TI900 Triboindenter.
(100)
deformation temperature or lower strain rates [19,20]. However, the formation of nanograins in existing SNC techniques mainly originates from the compression strain at high strain rates and/or cryogenic temperatures. Accordingly, if large shear strain at high strain rates and low temperature are introduced into the SNC process, both the processing method and the grain refining mechanism will undoubtedly be interesting issues. Zirconium (Zr) with an hcp structure has important applications in nuclear industry, aviation and surgical implant due to its resistance to corrosion and irradiation, high melting point and good biocompatibility [21]. In the present work, we chose Zr as a model metal and developed a novel and effective technique, namely surface circulation rolling treatment (SCRT), to fabricate a large-area nanostructured surface on Zr plate at cryogenic temperatures. The microstructure evolutions and hardness of nanocrystalline Zr are investigated in details.
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II
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Fig. 3. Cross-sectional SEM image of the zirconium sample after SCRT.
3.2. Scanning electron microscopy observations The cross-sectional SEM image of the SCRT sample is shown in Fig. 3. It is obvious that after SCRT, three distinct regions can be recognized: (I) the top surface region (distanced the surface to about 300 mm) which experienced the severe plastic deformation. In this region, the original grain boundaries could hardly be identified under the SEM, indicating surface grain refinement induced by SCRT; (II) the transition region (distanced the surface from about 300 to 600 mm) which was affected by smaller plastic deformation. Comparing to the top surface region, part of the grain boundaries were clearly visible and the grains were elongated to a certain extent; (III) the strain-free coarse-grained region (distanced the surface to about 600 mm) which was not affected during SCRT process. Accordingly, the depth of plastic deformation layer can be deduced to exceed 600 mm, which is higher than that obtained by means of SMAT [7,8] and SMGT [11]. To take insight into the microstructure evolutions of the deformed layers, based on the SEM image, the TEM observations of the microstructures at 400, 250, 100 and 30 mm depths from the surface were carried out. 3.3. Quantitative analysis on microstructure evolution As shown in Fig. 4(a), the microstructure of the layer at the depth of approximately 400 mm from the surface displays parallel deformation bands, which are several hundreds of nanometers in width and a few micrometers in length. These deformation bands divided the Zr coarse grains into finer plates. Meanwhile, many slip systems located in the deformation bands were activated. High-density dislocations and interactions occurred to form dislocation cells and pile up at the boundaries because the boundaries acted as obstacles for dislocation slips [24]. Deduction can be made that deformation bands dominate the plastic deformation at the low strain. It is also consistent with previous report [8,25]. Fig. 4(b) shows the microstructure of the layer at the depth of about 250 mm. Since the relative atomic movement is limited in the deformation bands, the gross deformation processed by deformation bands is quite small, and most of the plastic flow is due to movement of dislocations [26,27]. Hence, with further increase in the strain, the dislocation slips and interactions became the dominant deformation mode instead of deformation bands. The dislocation rearrangement led to the formation of dislocation cells that minimized the total system energy. The similar results are also observed in other hcp metals, such as Ti [17] and Mg [18]. In addition, the intersection between the deformation bands reported in the deformed layers of SMAT [8] was not evident. Compared with SMAT and SMGT, in which only a local small-area shear deformation occurred, the SCRT process produced large-area shear deformation. This deformation mode may be more beneficial to the activation of dislocation slips. As a result, a large number of dislocations emerged within the deformation bands, which further subdivided the plates into small blocks. With increasing strain, the accumulation and
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rearrangement of dislocations resulted in direct transformation of small blocks into the dislocation cells without the intersection of the deformation bands. The microstructure of the layer at the depth of about 100 mm is shown in Fig. 4(c). Dislocation cells continuously transformed into submicronic grains with the increase in the applied strain. The grain size decreased from more than 200 nm (marked A) to less than 100 nm (marked B). It was also found that dislocations (marked with arrows) pile up inside the grains. Fig. 4(d) shows the microstructure of the layer at the depth of about 30 mm. The grain sizes were less than 100 nm. DDWs (marked with arrows) occurred because of dislocation accumulation and rearrangement. The corresponding SAED pattern with elongated and discrete diffraction arcs indicated that the grain orientations were random, and the submicronic grains shown in Fig. 4(c) started to transform into nanograins. It is believed that the increase of strain is responsible for the subdivision of the subgrains into nanograins [8]. Fig. 4(e) shows a TEM bright-field image taken from the topmost surface of the specimen after SCRT. Under very high strain, the DDWs transformed into grain boundaries, and the submicronic grains were subdivided into equiaxial nanograins. The inset of Fig. 4(e) shows the corresponding SAED pattern with a series of rings, which indicates the random crystallographic orientations of the nanograins. The histogram of the grain-size distribution obtained from a large number of dark-field TEM images was characterized by normal logarithmic distribution with a narrow-size distribution ranging from 4 nm to 30 nm (as shown in Fig. 4f). The average grain size was approximately 8 nm, significantly smaller than that found by X-ray diffraction (XRD) analysis (42 nm). The discrepancy can be ascribed to the fact that the XRD structure information averaged approximately 5 mm thick at the top surface layer [7], as well as to surface roughness. Actually, when the grain sizes are less than a critical value, the dislocation multiplication rate is balanced by the annihilation rate [7]. As a result, dislocations will not accumulate in the grains, and the nanograins become free of dislocations, i.e., the increase in strains could not significantly reduce the nanograin size [7,8]. With increasing depth from the treated surface, grain sizes gradually increase. The distribution of the average structure size determined from a large number of TEM images with the corresponding distance from the treated surface is shown in Fig. 5. An increasing trend is obvious for the grain sizes with increasing depth. A gradient structure is formed with a continuously increasing grain size from nano-scale to the submicron and micron scales in the treated surface layer. Within the top 80 mm thick layer, the grain sizes are less than 100 nm. Hence, a nanostructured layer about 80 mm thick is achieved in the sample by using SCRT. 3.4. Nanoindentation hardness measurements The distributions of hardness and elastic modulus across the treated surface layer are shown in Fig. 6(a). The hardness increased from approximately 2.86 GPa (or 291.84 in the Hv scale) in the coarse-grained matrix to approximately 6.0 GPa at the topmost surface, in accordance with the classic Hall–Petch (H–P) relationship and consistent with the results in references [28] and [29]. The hardness of the top surface layer is about 2 times higher than that of the coarse-grained matrix. We believe that both the intrinsic high hardness and the nano-grains of the surface layer contribute to the observed high hardness of the surface layer of the SCRT specimen [30]. However, this hardness variation trend is much slower than that in previous reports [11,30], which could be more conducive to the performance of the material surface. On the other hand, the elastic modulus maintained a constant
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Fig. 4. TEM micrographs of the layer at the depth of about (a) 400 mm, (b) 250 mm, (c) 100 mm (inset: corresponding SAED pattern) and (d) 30 mm (inset: corresponding SAED pattern) and the top surface layer: (e) bright-field image (inset: corresponding SAED pattern) and (f) dark-field image (inset: histogram of grain size distribution derived from a total of 963 grains).
value of approximately 53.63 GPa over the entire deep range, indicating that the elastic modulus weakly depends on the grain size, also consistent with previous reports [11,31]. Fig. 6b presents the H–P relationship between the nanoindentation hardness, H, and the average grain size, d, revealing that the H–d 1/2 relationship is not linear, but is concave towards to the d 1/2
axis. The hardness as a function of the inverse of the square root of the average grain size is linear at locations 20 mm away from the treated surface. However, within a range of 20 mm from the treated surface, the nanohardness data deviate from the straight line and bend towards the d 1/2 axis. The non-linear H–d 1/2 relationship may due to the change in the deformation mechanism as the average
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grain boundaries. It is consistent with the TEM results above. For small grain sizes (i.e. do17 nm), the dislocations existing in the grains are rare, or even only one. So the H–P relationship is not applicable. Furthermore, the grain boundary regions of nanocrystalline materials will appear relaxation process under the stress, the hardness of materials decreases. The hardness has been dictated by both dislocation interactions and grain-boundary sliding acting simultaneously [32–35]. The transition of the deformation mechanism occurs in the intermediate size of the grain sizes ranging from 17 nm to 25 nm.
100000 10000 1000 250 200 150 100
4. Conclusions
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Distance from surface (µm) Fig. 5. Variations of the average structure size with distance from treated surface.
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A novel and efficient SNC technique, namely SCRT, has been developed to fabricate a nanostructured surface layer on commercial pure zirconium. The thickness of the deformation layer exceeded 600 mm. The average grain sizes varied from about 8 nm in the topmost surface to micrometers in the coarse grained matrix. Meanwhile, a nanostructured layer about 80 mm thick was achieved in the sample. The occurrence of deformation bands was responsible for the grain refinement during the initial stage of deformation. With increasing strain, dislocation cells were formed in the deformation bands and finally transformed into nanograins. The SCRT layer hardness exhibited a gradient variation from approximately 6.0 GPa at the top surface to 2.86 GPa in the coarse-grained matrix, whereas the elastic modulus retained a constant value of approximately 53.63 GPa throughout the entire deep range. However, the Hall–Petch relationship between hardness and average grain size was not linear due to the change of the deformation mechanism from dislocation pile-ups to grainboundary sliding as the grain size became smaller. The SCRT technique can easily manufacture large-area nanostructured layer on the metal plate by a simple process. The improved surface performance is worth to expect.
Distance from surface (μm) Acknowledgments
Distance from surface (μm) 600 400 300 10050
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This work was supported in part by the National Basic Research Program of China (Grant no. 2010CB731606) and in part by the National Science Foundation for Distinguished Young Scholars (Grant no. 50925522).
6.5 60 2 280 120 76 6.0
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References 5.0
[1] [2] [3] [4]
4.5 4.0
[5] 3.5
[6] [7]
3.0 2.5 0
2
4
6
8
10
12
14
Inverse of the square root of the average grain size (μm-1/2) Fig. 6. (a) Variations of hardness and elastic modulus with distance from treated surface and (b) the hardness, H, as a function of the inverse of the square root of the average grain size, d 1/2 Fig. 2.
grain size becomes smaller. For large grain sizes (i.e. d & 25 nm), the hardness is mainly controlled by the increase in the internal stress due to the dislocation pile-ups which results from the presence of
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556. H. Gleiter, Acta Mater. 48 (2000) 1–29. W.L. Li, N.R. Tao, Z. Han, K. Lu, Wear 274–275 (2012) 306–312. K.S. Kumar, H. Van Swygenhoven, S. Suresh, Acta Mater. 51 (2003) 5743–5774. M. Dao, L. Lu, R.J. Asaro, J.T.M. De Hosson, E. Ma, Acta Mater. 55 (2007) 4041–4065. K. Lu, J. Lu, J. Mater. Sci. Technol. 15 (1999) 193–197. N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, Acta Mater. 50 (2002) 4603–4616. L. Zhang, Y. Han, J. Lu, Nanotechnology 19 (2008) 165706. G. Liu, S.C. Wang, X.F. Lou, J. Lu, K. Lu, Scr. Mater. 44 (2001) 1791–1795. G. Liu, J. Lu, K. Lu, Mater. Sci. Eng. A 286 (2000) 91–95. W.L. Li, N.R. Tao, K. Lu, Scr. Mater. 59 (2008) 546–549. T.H. Fang, W.L. Li, N.R. Tao, K. Lu, Science 331 (2011) 1587–1590. R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Prog. Mater. Sci. 45 (2000) 103–189. K. Lu, J. Lu, Mater. Sci. Eng. A 375–377 (2004) 38–45. Y.B. Wang, M. Louie, Y. Cao, X.Z. Liao, H.J. Li, S.P. Ringer, Y.T. Zhu, Scr. Mater. 62 (2010) 214–217. K. Wang, N.R. Tao, G. Liu, J. Lu, K. Lu, Acta Mater. 54 (2006) 5281–5291. K.Y. Zhu, A. Vassel, F. Brisset, K. Lu, J. Lu, Acta Mater. 52 (2004) 4101–4110. H.Q. Sun, Y.N. Shi, M.X. Zhang, K. Lu, Acta Mater. 55 (2007) 975–982. Q. Wei, Z.L. Pan, X.L. Wu, B.E. Schuster, L.J. Kecskes, R.Z. Valiev, Acta Mater. 59 (2011) 2423–2436.
32
C. Yuan et al. / Materials Science & Engineering A 565 (2013) 27–32
[20] G.Q. Zhao, S.B. Xu, Y.G. Luan, Y.J. Guan, N. Lun, X.F. Ren, Mater. Sci. Eng. A 437 (2006) 281–292. [21] D.F. Guo, M. Li, Y.D. Shi, Z.B. Zhang, H.T. Zhang, X.M. Liu, B.N. Wei, X.Y. Zhang, Mater. Des 34 (2012) 275–278. [22] M. Wen, G. Liu, J.F. Gu, W.M. Guan, J. Lu., Surf. Coat. Technol 202 (2008) 4728–4733. [23] H.P. Klug, L.E. Alexander, X-Ray Diffraction Procedures for Polycrystalline and Amorphous Materials, Wiley, New York, 1974, p. 661. [24] L. Lu, Y.F. Shen, X.H. Chen, L.H. Qian, K. Lu, Science 304 (2004) 422–426. [25] W.S. Choi, H.S. Ryoo, S.K. Hwang, M.H. Kim, Metall. Mater. Trans. A 33 (2002) 973–980. [26] T.H. Courtney, Mechanical Behavior of Materials, McGraw-Hill, New York, 1990, p. 309.
[27] R.E. Reed-Hill, Physical Metallurgy Principles, Affiliated East, West Press, New Delhi, 1973. [28] E.O. Hall, Proc. Phys. Soc. London Sect B 64 (1951) 747–753. [29] N.J. Petch, J. Iron Steel Inst. London 174 (1953) 25. [30] P. Jiang, Q. Wei, Y.S. Hong, J. Lu, X.L. Wu, Surf. Coat. Technol 202 (2007) 583–589. [31] T.D. Shen, C.C. Koch, T.Y. Tsui, G.M. Pharr, J. Mater. Res. 10 (1995) 2892–2896. [32] L.L. Shaw, A.L. Ortiz, J.C. Villegas, Scr. Mater. 58 (2008) 951–954. [33] K.S. Kumar, H. Van Swygenhoven, S. Suresh, Acta Mater. 51 (2003) 5743–5774. [34] K.A. Padmanabhan, G.P. Dinda, H. Hahn, H. Gleiter, Mater. Sci. Eng. A (2007) 462–468452-3 (2007) 462–468. [35] C.S. Pande, K.P. Cooper, Prog. Mater. Sci. 54 (2009) 689–706.