Effects of Deep Cryogenic Treatment on the Microstructure and Mechanical Properties of Commercial Pure Zirconium 2015 Journal of Alloys and Compounds | Plasticity (Physics) | Hardness

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  Effects of deep cryogenic treatment on the microstructure andmechanical properties of commercial pure zirconium Chao Yuan, Yunpeng Wang, Deli Sang, Yijun Li, Lei Jing, Ruidong Fu ⇑ , Xiangyi Zhang State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, Hebei 066004, PR ChinaCollege of Materials Science and Engineering, Yanshan University, Qinhuangdao, Hebei 066004, PR China a r t i c l e i n f o  Article history: Received 25 April 2014Received in revised form 3 August 2014Accepted 26 August 2014Available online 16 September 2014 Keywords: Metals and alloysDeep cryogenic treatmentMicrostructureMechanical propertiesZirconium a b s t r a c t The effects of deep cryogenic treatment (DCT) on the microstructure and mechanical properties of com-mercial pure zirconium were investigated. Experimental results indicated that DCT induced a change ingrain orientation and improved internal stress, which in turn increased dislocation density that led toimproved hardness. Hardness in basal planes was found to be significantly larger than that in prismplanes. Moreover, strength was enhanced in DCT-treated zirconium and the ductility was comparableto that of as-annealed zirconium. This phenomenon was due to the increase in dislocation density andthe good ductility resulting from the motion of pre-existing dislocations and specific dislocation config-urations. DCT led to the transformation of tensile fracture mode from mixed-rupture characteristics of quasi-cleavage and dimples to quasi-cleavage, thereby increasing compatible deformation capabilities.The possible mechanisms underlying microstructural modification, tensile strength, and hardnessimprovement were discussed.   2014 Elsevier B.V. All rights reserved. 1. Introduction Cryogenic treatment is a very old process that is widely used forhigh-precision components. Shallow cryogenic treatment is set atlow temperatures (about   80   C), whereas deep cryogenic treat-ment (DCT) is set at near-liquid-nitrogen temperatures (about  196   C). DCT improves certain properties beyond the enhance-ment obtained by normal cold treatment [1–3]. DCT is differentfrom those methods of severe plastic deformation (SPD) processingat cryogenic conditions, such as cryogenic rolling, pinning andmilling. Instead of nanostructured and ultrafine-grains in SPD-pro-cessedmetals,thevariationofcrystaldefectand phasetransforma-tion may play more important role in increasing of the propertiesof DCT-treated metals. Over the past few decades, the effects of various cryogenic treatments on the performance of steel havereceived considerable interest. DCT reportedly increases the nor-mal temperature strength and hardness of steels [4,5], providesdimensional stability or microstructural stability [6], and improveswear [7–9] and fatigue resistance [10,11]. The improvement in mechanical properties can be ascribed to the complete transforma-tion of retained austenite into martensite, precipitation of fine dis-persed carbides, and removal of residual stresses [6,7,12]. Previousstudies have shown that the hardness and abrasion resistance of DCT-treated samples evidently improve because of transformationof abundant retained austenite to martensite, secondary carbideprecipitation [13], and the precipitation of nanosized  g -carbidesin primary martensite [14]. In addition, DCT combined with tem-pering treatment enables the enhancement of fatigue propertiesfor the precipitation of fine carbides and the reduction in compres-sive residual stress [15].Compared with studies on ferrous metals, research on the effectof DCT on nonferrous metals such as Mg, Al, and Ti alloys are lim-ited. Asl et al. [2] found that after DCT, tiny laminar  b  phase parti-cles almost dissolve, and the coarse divorced eutectic  b  phaseextends into the neighboring matrix. As a result, the mechanicalproperties of AZ91 Mg alloy significantly improve. Jiang et al.[16] found that DCT induces the refinement of grains to approxi-mately 0.1–3.0 l m, thereby improving the strength and elongationof 3102 Al alloy foil.As aforementioned, evident changes in microstructure such asphase transformation or grain refinement have been verified forDCT-treated alloys. However, changes in microstructure andmechanical properties that can be introduced by DCT if no phasetransformation occurs are worth investigating. In this study,coarse-grained zirconium (Zr) with a hexagonal close packed(hcp) structure and higher microstructure stability at cryogenic http://dx.doi.org/10.1016/j.jallcom.2014.08.2010925-8388/   2014 Elsevier B.V. All rights reserved. ⇑ Corresponding author at: State Key Laboratory of Metastable Materials Scienceand Technology, Yanshan University, Qinhuangdao, Hebei 066004, PR China. Tel.:+86 335 858 7046; fax: +86 335 807 4545. E-mail address:  rdfu@ysu.edu.cn (R. Fu). Journal of Alloys and Compounds 619 (2015) 513–519 Contents lists available at ScienceDirect  Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom  temperature was chosen as a model metal because of the limiteddeformation capability that results in easily saturated dislocationdensity [17]. Some unusual changes in mechanical properties werefound in DCT-treated Zr, and the possible mechanisms in terms of microstructural modification were discussed. 2. Experimental The material used in this study was commercial pure Zr plate (chemicalcomposition summarized in Table 1). The plate was annealed in a vacuum at1123 K for 4 h to obtain a homogeneous polycrystalline structure. The average grainsize of the as-annealed sample was approximately 30–50 l m. The samples cut fromthe as-annealed plate were placed in a liquid nitrogen environment at 77 K for 24 h.Afterwards, the samples were taken out and cooled to room temperature.X-ray diffraction (XRD) patterns of the as-annealed and DCT-treated sampleswere measured using a Rigaku D/max 2500 X-ray diffractometer (18 kW) with CuK a  radiation (wavelength  k  = 1.54 Å) in continuous-scanning mode over 2 h  = 30–80   at a step of 0.02  . Differential scanning calorimetry (DSC) was conducted on aDiamond DSC (PerkinElmer Inc., UK). For dynamic scans, samples were subjectedto temperatures ranging from 20   C to   160   C at a cooling rate of 20   C/min.Microstructure features before and after DCT were observed by optical microscopy(OM), electron backscatter diffraction (EBSD), and transmission electron micros-copy (TEM). EBSD analysis was performed using a HITACHI S-4800 scanning elec-tron microscopy (SEM) system, whereas TEM observations were carried out witha JEOL-2010 TEM system at a voltage of 200 kV.A FM-ARS9000 Vickers microhardness tester was used to measure the hardnessof samples. Indentations were carried out with a given load of 200 g and a dwelltime of 10 s. Hardness distributions over a square area of 4.5    4.5 mm 2 , with aspace between adjacent indentations of 0.5 mm were obtained. The tensile sampleswere machined with a gauge length of 30 mm and a cross-section of 10    1.5 mm 2 .Four repeated tensile tests were performed on a MTS test system at a strain rate of 1    10  4 s  1 at room temperature under extensometer-measured strain control.Fracture surfaces after the tensile test were observed by SEM, and microstructuresnear the fracture surfaces were observed by TEM. 3. Results and discussion  3.1. XRD and DSC analyses Fig. 1(a) shows the XRD patterns of as-annealed and DCT-treated samples. No phase transformation was observed in DCT-treated Zr compared with as-annealed Zr. However, the intensitiesof the diffraction peaks of DCT-treated Zr were slightly enhanced,which can be attributed to the variation in lattice constants causedby DCT. Changes in lattice constants were also observed in DCT-treated Mg [18,19]. To further confirm the microstructure stabilityofZratcryogenictemperature,acryogenicDSCcurvewasobtainedand is shown in Fig. 1(b). The smooth feature of the DSC curve indi-cated that no transformation occurred during DCT.  3.2. EBSD analysis To gain deep insight into microstructural changes, EBSD analy-sis was performed.Inthis analysis,orientation imagingmicroscopy(OIM) maps, inverse pole figure (IPF) maps, and boundary misori-entation angle distributions were obtained before and after DCT,as shown in Fig. 2. Fig. 2(a) and (b) shows the OIM maps of Zr before and after DCT, with high-angle (HAGBs; grain boundarymisorientations P 15  ) and low-angle (LAGBs; grain boundarymisorientations < 15  ) grain boundaries depicted by black andwhite lines, respectively. The crystallographic directions corre-sponding to various colors can be inferred from the IPF triangleshown at the bottom right corner of  Fig. 2(a) and (b). The micro-structure was characterized by equiaxed grains and a higherfraction of HAGBs. No phase transformation and precipitationoccurred during DCT, consistent with the XRD and DSC results.Fig. 2(c) and (d) shows the IPF maps of Zr before and after DCT,respectively. The maximum intensity indicated the number of ran-dom orientations. A strong preference was observed toward the(0001) orientation, which was corroborated by the OIM maps inFig. 2(a) and (b). Furthermore, the intensity of prism planes waslower and the grain orientations were much closer to the (0001)basal plane orientation after DCT. This finding can be attributedto the ordered arrangement of atoms resulting from the smallchanges in lattice constants [18,19]. The misorientation angle dis-tributions in Fig. 2(e) and (f) shows that the crystallite boundarieswere mainly high angle in nature. The HAGB fractions of Zr beforeand after DCTwere about 81.2%and 85.3%of the total grain bound-ary length, respectively. Thus, DCT improved the formation of HAGBs. The increase in HAGB fractions may be related to theextrinsic dislocations formed during DCT (Fig. 3). The extrinsic dis-locations were probably easier to react with the dislocations inLAGBs [20,21]. Therefore, the accumulation and rearrangement of dislocations resulted in the transformation of LAGBs to HAGBs,leading to increased HAGB and decreased LAGB fractions.  3.3. TEM observations The TEM images of the samples before and after DCT are pre-sented in Fig. 3. The dislocation density of the DCT-treated samplewas apparently higher than that of the as-annealed sample. Whenthe temperature dropped from room temperature to cryogenic  Table 1 Chemical composition (in wt%) of the investigated pure Zr. Fe + Cr C N H O Hf Zr0.2 0.05 0.01 0.005 0.16 4.5 Balance Fig. 1.  XRD patterns of as-annealed and DCT-treated samples (a), and the cryogenicDSC curve of Zr (b).514  C. Yuan et al./Journal of Alloys and Compounds 619 (2015) 513–519  temperature, the volume of the material contracted. This volumecontraction released great compression deformation energy thatserved as the driving force for the formation and movement of dis-locations [18]. The improvement in dislocation density played asignificant role in enhancing mechanical properties and deforma-tion behavior, as discussed in the following sections.  3.4. Hardness measurements To better understand the variation in hardness of Zr, contourmaps of the spatial distribution of the Vickers hardness ( H  v) overa 4.5    4.5 mm 2 square area before and after DCT are shown inFig. 4(a) and (b), respectively. The hardness of as-annealed Zrsubstantially varied from approximately 140 – 190 H  v, with an averagevalue of 159.5  H  v. By contrast, the hardness of DCT-treated Zr gen-erally changed from approximately 160 – 220  H  v, with an averagevalue of 197.5  H  v. Evidently, the average hardness significantlyincreased after DCT. The difference between the two average hard-ness values reached 23.8%. Although the spatial distribution of hardness was not uniform from one area to another, the increasein hardness of the DCT-treated sample was a holistic rather thanlocalized phenomenon, as can also be confirmed by the hardnessevolution characterized by the color maps in Fig. 4. The inhomoge-neity of hardness distribution may be closely related to the grainorientation. An in situ comparison between the OIM map andOM image of hardness indentation of DCT-treated Zr is presented Fig. 2.  Representative OIM maps of (a) as-annealed and (b) DCT-treated samples. IPF maps of (c) as-annealed and (d) DCT-treated samples. Distribution of boundarymisorientation angle of (e) as-annealed and (f) DCT-treated samples. C. Yuan et al./Journal of Alloys and Compounds 619 (2015) 513–519  515  in Fig. 5. The letters in Fig. 5(a) correspond to those in Fig. 5(b), and grain M indicates the referencegrain. The hardness values of grainsA–D in Fig. 5(b) were 211.65, 259.54, 199.35, and 233.57  H  v,respectively. The hardness values varied with the orientationbetween basal and prism planes. Planes close to the (0001) basalplane orientation were significantly harder than planes close tothe prism  ð 10  10 Þ  and  ð 2  1  10 Þ  planes. The hardness differenceamongdifferentgrainorientationsreached upto 30.2%.Thedepen-dency of hardness values on orientation has also been observed inprevious reports [22–24].For hcp Zr, the axial ratio  c/a  was 1.593, which was less thanthat of the ideal value (i.e., <1.633). Consequently, the lattice resis-tance for prism planes was lower than that for basal or pyramidalplanes; thus, slip preferred to occur on prism planes [25]. In otherwords, plastic deformation more easily occurred on prism planesthan on basal planes. Consequently, the hardness in basal planeswas much larger than that in prism planes. Therefore, the overallimprovement of hardness may be closely related to both changesin grain orientation and dislocation density. After DCT, the grainorientation tended to be more consistent, resulting in orderedstrengthening [18]. Meanwhile, the increase in dislocation densityincreased the force to overcome the resistance of dislocationmotion. Thus, the improvement in plastic deformation resistanceresulted in increased hardness.  3.5. Tensile properties The tensile engineering and true stress–strain curves of as-annealed and DCT-treated Zr are presented in Fig. 6. Foras-annealed Zr, the yield strength (0.2% offset,  r 0.2 ) and ultimatetensile strength ( r UTS ) were 288 and 393 MPa, respectively. Theelongation to fracture reached 28.7%. Correspondingly,  r 0.2  and r UTS  of DCT-treated Zr increased by about 60% and 40%, respec-tively, compared with as-annealed Zr. Meanwhile, the elongationto fracture reached 26.5%, which was similar to that of as-annealedZr. This finding indicated that DCT was a better route to improvingthe strength and keeping a good ductility of Zr.As depicted in Fig. 6(b), DCT-treated Zr also exhibited highstrain hardening capabilities relative to as-annealed Zr. The strainhardening exponent ( n ) is defined by [26] r ¼  K  e n ð 1 Þ where  r  is the true stress,  e  is the true strain, and  K   is a constant.The  n  values for as-annealed and DCT-treated Zr were 0.09 and0.11,respectively. The enhancementin strainhardeningcapabilitiesof DCT-treated Zr was further confirmed by an improved strainhardening rate ( H ) defined by [26] H  ¼  1 r @  r @  e    ð 2 Þ Variations in  H  with  e  before and after DCT are displayed inFig. 7. The variation trends were consistent with continuous strainhardening to significant strains, but thestrain hardening rate of theDCT-treated sample was larger than that of the as-annealedsample. Fig. 3.  TEM images of (a) as-annealed and (b) DCT-treated samples. Fig. 4.  Spatial distributions of the hardness values of Zr: (a) before and (b) afterDCT.516  C. Yuan et al./Journal of Alloys and Compounds 619 (2015) 513–519
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