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Microstructure evolution and mechanical properties of nanocrystallinezirconium 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 Chinab College of Materials Science and Engineering, Yanshan University, Qinhuangdao, Hebei 066004, PR Chinac School of Mechanical & Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore

a r t i c l e i n f o

Article history:

Received 11 October 2012

Received in revised form

13 November 2012

Accepted 15 November 2012Available online 12 December 2012

Keywords:

Nanocrystalline materials

Microstructure

Nanoindentation hardness

Surface circulation rolling treatment (SCRT)

Zirconium

a b s t r a c t

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 grain-

boundary sliding as the grain size becomes smaller.

& 2012 Elsevier B.V. All rights reserved.

1. Introduction

Nanocrystalline materials exhibit many excellent properties,such as excellent hardness and strength, outstanding tribologicaland magnetic properties, and improved toughness, comparedwith conventional polycrystalline materials [1–5]. Consequently,these materials have been extensively researched in the pastdecades. However, difficulties in preparing bulk nanocrystallinematerials limit their engineering applications. For metals, mostfailures originate from the surface, whose surface microstructuresand properties are very sensitive. Hence, global performances ofmetallic materials can be improved by surface modification [6].Lu and Lu [6] first proposed the concept of surface nanocrystalli-zation (SNC) of metallic materials. Since then, fabrication ofnanostructured layer on a bulk material surface has receivedincreasing attention. Various mechanical processing techniqueshave been successfully used to fabricate nanostructures onmetallic material surfaces. Such techniques include surfacemechanical attrition treatment (SMAT) [7,8], high-energy shotpeening [9], ultrasonic shot peening [10], surface mechanicalgrinding treatment (SMGT) [11,12], and so on. Meanwhile, theperformance of nanostructured layers has also been investigated.

For example, an ideal gradient nano-micro-grained architecturewas synthesized in a bulk coarse-grained Cu substrate. The yieldstrength of the gradient nano-/coarse-grained structure was twicethat of the coarse-grained substrate. The uniform elongationbetween these substrates was exactly the same [12]. These resultswere desirable and profoundly influenced subsequent studies.

The grain refinement mechanism in the SNC process is similarto that in severe plastic deformation (SPD) techniques, exten-sively used to refine grains in various metals and alloys over thepast decades [13,14]. The mechanism of SPD-induced grainrefinement can usually be attributed to phase transformation,dislocation activities, and/or deformation twinning. Generally,crystal structure, stacking fault energy (SFE), and initial grain sizewere considered as the intrinsic factors that affect the aboveprocesses [15]. For example, for bcc Fe with a high SFE, theformation of dislocation walls and dislocation tangles resulted inthe grain refinement [7]; for fcc Cu with a medium SFE, grainswere refined via formation of dislocation cells, dislocation wallsand twins [16]; while for hcp Ti with a high SFE and Mg with amedium SFE, the grain refinement process involved the formationof 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 asthe extrinsic factors that govern the grain refinement process. Forexample, the two most important SPD techniques, namely equal-channel angular pressing and high-pressure torsion, mainly drawlarge shear strain from the deformation process at higher

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Materials Science & Engineering A

0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved.

http://dx.doi.org/10.1016/j.msea.2012.11.092

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: þ86 335 807 4545.

E-mail address: rdfu@ysu.edu.cn (R. Fu).

Materials Science & Engineering A 565 (2013) 27–32

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deformation temperature or lower strain rates [19,20]. However,the formation of nanograins in existing SNC techniques mainlyoriginates from the compression strain at high strain rates and/orcryogenic temperatures. Accordingly, if large shear strain at highstrain rates and low temperature are introduced into the SNCprocess, both the processing method and the grain refiningmechanism will undoubtedly be interesting issues.

Zirconium (Zr) with an hcp structure has important applica-tions in nuclear industry, aviation and surgical implant due to itsresistance to corrosion and irradiation, high melting point andgood biocompatibility [21]. In the present work, we chose Zr as amodel metal and developed a novel and effective technique,namely surface circulation rolling treatment (SCRT), to fabricatea large-area nanostructured surface on Zr plate at cryogenictemperatures. The microstructure evolutions and hardness ofnanocrystalline Zr are investigated in details.

2. Experimental

A commercial pure zirconium (Zr 702) plate (90 mm�65 mm�3 mm) was used for the SCRT experiment. The platewas annealed at 1223 K for 4 h to obtain a coarse-grainedpolycrystalline structure. The average grain size of theas-annealed sample is about 60 mm.

Fig. 1 shows a schematic illustration of the SCRT set-up used inthe present work. A cylindrical tool with a curved surface at thetip (Fig. 1a) rotates at rotational speed o and maintains apenetration depth d into the sample surface. The sample (or tool)moves at a travel speed u along the horizontal direction. Theplastic strain and depth of the deformation layer depend on thepenetration depth and the tool tip curvature radius r (Fig. 1b). Inorder to refine the original coarse grains into nanograins, effectivecooling of the processed material and the tool is necessaryduring SCRT.

Before the SCRT the plate sample and the tool were immersedinto liquid nitrogen for 5 min in order to ensure effectivepre-cooling. During the SCRT the plate sample was soaked inliquid nitrogen all the time. The SCRT processing parameters aredescribed as follows: o¼400 rpm, u¼100 mm min�1, d¼0.1 mm,r¼40 mm. In order to effectively refine the grains in the surfacelayer and increase the thickness of the nanostructured layer, theSCRT process was repeated three times with the same processingparameters. A total reduction of 0.2 mm and a strain rate of about104–105 s�1 were achieved in the SCRT Zr.

The cross-sectional observations of the treated Zr sample werepreformed on a HITACHI S-4800 scanning electron microscope(SEM). The microstructure features in the surface layer werecharacterized by using a JEOL-2010 transmission electron micro-scope (TEM) at an operating voltage of 200 kV. The cross-sectionalhardness and elastic modulus profile of the SCRT surface layerwere measured by using a TI900 Triboindenter.

3. Results and discussion

3.1. X-ray diffraction analysis

Fig. 2 shows the XRD patterns of the surface layer of theSCRTed and un-SCRTed samples. The Miller indices of the reflect-ing planes corresponding to each peak are given in the figure. It isnoted that SCRT does not result in any phase transformation of Zr;however, evident Bragg-diffractional peak broadening exists inthe SCRTed Zr, which is caused by the grain refinement and anincrease in the atomic-level microstrain [22]. The average grainsize calculated by the Scherrer and Wilson method [23] was about42 nm, and the microstrain was about 0.18%.

ω

υ

dr

Tool tip

Deformed layerSample

Fig. 1. Schematic illustrations of (a) the tool tip outline and (b) the SCRT process.

30 35 40 45 50 55 60 65 70 75 80

(202

)

(004

)

(201

)(1

12)

(200

)

(103

)

(110

)(102

)

(002

)(1

01)

un-SCRT SCRT

Inte

nsity

2-theta (degree)

(100

)

Fig. 2. XRD patterns of the SCRTed and un-SCRTed samples.

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3.2. Scanning electron microscopy observations

The cross-sectional SEM image of the SCRT sample is shown inFig. 3. It is obvious that after SCRT, three distinct regions can berecognized: (I) the top surface region (distanced the surface toabout 300 mm) which experienced the severe plastic deformation.In this region, the original grain boundaries could hardly beidentified under the SEM, indicating surface grain refinementinduced by SCRT; (II) the transition region (distanced the surfacefrom about 300 to 600 mm) which was affected by smaller plasticdeformation. Comparing to the top surface region, part of thegrain boundaries were clearly visible and the grains were elon-gated to a certain extent; (III) the strain-free coarse-grainedregion (distanced the surface to about 600 mm) which was notaffected during SCRT process. Accordingly, the depth of plasticdeformation layer can be deduced to exceed 600 mm, which ishigher than that obtained by means of SMAT [7,8] and SMGT [11].To take insight into the microstructure evolutions of thedeformed layers, based on the SEM image, the TEM observationsof the microstructures at 400, 250, 100 and 30 mm depths fromthe surface were carried out.

3.3. Quantitative analysis on microstructure evolution

As shown in Fig. 4(a), the microstructure of the layer at thedepth of approximately 400 mm from the surface displays paralleldeformation bands, which are several hundreds of nanometers inwidth and a few micrometers in length. These deformation bandsdivided the Zr coarse grains into finer plates. Meanwhile, manyslip systems located in the deformation bands were activated.High-density dislocations and interactions occurred to formdislocation cells and pile up at the boundaries because theboundaries acted as obstacles for dislocation slips [24]. Deductioncan be made that deformation bands dominate the plasticdeformation at the low strain. It is also consistent with previousreport [8,25]. Fig. 4(b) shows the microstructure of the layer atthe depth of about 250 mm. Since the relative atomic movement islimited in the deformation bands, the gross deformation pro-cessed by deformation bands is quite small, and most of theplastic flow is due to movement of dislocations [26,27]. Hence,with further increase in the strain, the dislocation slips andinteractions became the dominant deformation mode instead ofdeformation bands. The dislocation rearrangement led to theformation of dislocation cells that minimized the total systemenergy. The similar results are also observed in other hcp metals,such as Ti [17] and Mg [18]. In addition, the intersection betweenthe deformation bands reported in the deformed layers of SMAT[8] was not evident. Compared with SMAT and SMGT, in whichonly a local small-area shear deformation occurred, the SCRTprocess produced large-area shear deformation. This deformationmode may be more beneficial to the activation of dislocation slips.As a result, a large number of dislocations emerged within thedeformation bands, which further subdivided the plates intosmall blocks. With increasing strain, the accumulation and

rearrangement of dislocations resulted in direct transformationof small blocks into the dislocation cells without the intersectionof the deformation bands.

The microstructure of the layer at the depth of about 100 mm isshown in Fig. 4(c). Dislocation cells continuously transformed intosubmicronic grains with the increase in the applied strain. Thegrain size decreased from more than 200 nm (marked A) to lessthan 100 nm (marked B). It was also found that dislocations(marked with arrows) pile up inside the grains. Fig. 4(d) showsthe microstructure of the layer at the depth of about 30 mm. Thegrain sizes were less than 100 nm. DDWs (marked with arrows)occurred because of dislocation accumulation and rearrangement.The corresponding SAED pattern with elongated and discretediffraction arcs indicated that the grain orientations wererandom, and the submicronic grains shown in Fig. 4(c) startedto transform into nanograins. It is believed that the increase ofstrain is responsible for the subdivision of the subgrains intonanograins [8].

Fig. 4(e) shows a TEM bright-field image taken from thetopmost surface of the specimen after SCRT. Under very highstrain, the DDWs transformed into grain boundaries, and thesubmicronic grains were subdivided into equiaxial nanograins.The inset of Fig. 4(e) shows the corresponding SAED pattern witha series of rings, which indicates the random crystallographicorientations of the nanograins. The histogram of the grain-sizedistribution obtained from a large number of dark-field TEMimages was characterized by normal logarithmic distributionwith a narrow-size distribution ranging from 4 nm to 30 nm (asshown 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 thatthe XRD structure information averaged approximately 5 mmthick at the top surface layer [7], as well as to surface roughness.Actually, when the grain sizes are less than a critical value, thedislocation multiplication rate is balanced by the annihilationrate [7]. As a result, dislocations will not accumulate in the grains,and the nanograins become free of dislocations, i.e., the increasein strains could not significantly reduce the nanograin size [7,8].

With increasing depth from the treated surface, grain sizesgradually increase. The distribution of the average structure sizedetermined from a large number of TEM images with thecorresponding distance from the treated surface is shown inFig. 5. An increasing trend is obvious for the grain sizes withincreasing depth. A gradient structure is formed with a continu-ously increasing grain size from nano-scale to the submicron andmicron scales in the treated surface layer. Within the top 80 mmthick layer, the grain sizes are less than 100 nm. Hence, ananostructured layer about 80 mm thick is achieved in the sampleby using SCRT.

3.4. Nanoindentation hardness measurements

The distributions of hardness and elastic modulus across thetreated surface layer are shown in Fig. 6(a). The hardnessincreased from approximately 2.86 GPa (or 291.84 in the Hv scale)in the coarse-grained matrix to approximately 6.0 GPa at thetopmost 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 higherthan that of the coarse-grained matrix. We believe that both theintrinsic high hardness and the nano-grains of the surface layercontribute to the observed high hardness of the surface layer ofthe SCRT specimen [30]. However, this hardness variation trend ismuch slower than that in previous reports [11,30], which could bemore conducive to the performance of the material surface.On the other hand, the elastic modulus maintained a constant

I II III

Trea

ted

surf

ace

100μm

Fig. 3. Cross-sectional SEM image of the zirconium sample after SCRT.

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value of approximately 53.63 GPa over the entire deep range,indicating that the elastic modulus weakly depends on the grainsize, also consistent with previous reports [11,31].

Fig. 6b presents the H–P relationship between the nanoindenta-tion hardness, H, and the average grain size, d, revealing that theH–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 ofthe average grain size is linear at locations 20 mm away from thetreated surface. However, within a range of 20 mm from the treatedsurface, the nanohardness data deviate from the straight line andbend towards the d�1/2 axis. The non-linear H–d�1/2 relationshipmay due to the change in the deformation mechanism as the average

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).

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grain size becomes smaller. For large grain sizes (i.e. d & 25 nm), thehardness is mainly controlled by the increase in the internal stressdue to the dislocation pile-ups which results from the presence of

grain boundaries. It is consistent with the TEM results above. Forsmall grain sizes (i.e. do17 nm), the dislocations existing in thegrains are rare, or even only one. So the H–P relationship is notapplicable. Furthermore, the grain boundary regions of nanocrys-talline materials will appear relaxation process under the stress,the hardness of materials decreases. The hardness has beendictated by both dislocation interactions and grain-boundarysliding acting simultaneously [32–35]. The transition of thedeformation mechanism occurs in the intermediate size of thegrain sizes ranging from 17 nm to 25 nm.

4. Conclusions

A novel and efficient SNC technique, namely SCRT, has beendeveloped to fabricate a nanostructured surface layer on com-mercial pure zirconium. The thickness of the deformation layerexceeded 600 mm. The average grain sizes varied from about 8 nmin the topmost surface to micrometers in the coarse grainedmatrix. Meanwhile, a nanostructured layer about 80 mm thickwas achieved in the sample. The occurrence of deformation bandswas responsible for the grain refinement during the initial stageof deformation. With increasing strain, dislocation cells wereformed in the deformation bands and finally transformed intonanograins. The SCRT layer hardness exhibited a gradient varia-tion from approximately 6.0 GPa at the top surface to 2.86 GPa inthe coarse-grained matrix, whereas the elastic modulus retained aconstant value of approximately 53.63 GPa throughout the entiredeep range. However, the Hall–Petch relationship between hard-ness and average grain size was not linear due to the change ofthe deformation mechanism from dislocation pile-ups to grain-boundary sliding as the grain size became smaller. The SCRTtechnique can easily manufacture large-area nanostructured layeron the metal plate by a simple process. The improved surfaceperformance is worth to expect.

Acknowledgments

This work was supported in part by the National BasicResearch Program of China (Grant no. 2010CB731606) and inpart by the National Science Foundation for Distinguished YoungScholars (Grant no. 50925522).

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0 100 200 300 400 500 600 700 800 900

50

100

150

200

250

1000

10000

100000

Ave

rage

str

uctu

re s

ize

(nm

)

Distance from surface (µm)

Fig. 5. Variations of the average structure size with distance from treated surface.

0 50 100 150 200 250 3000

1

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20

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(GPa

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Distance from surface (μm)

CG-Zr hardness

0 2 4 6 8 10 12 142.5

3.0

3.5

4.0

4.5

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6.0

6.5

Average grain size81317254076120280260

Distance from surface (μm)0510203050100300400

Har

dnes

s (G

Pa)

Inverse of the square root of the average grain size (μm-1/2)

600

µm nm

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.

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