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Investigation of microstructural evolution in a selective laser melted Ti6Al4V alloy induced by an ultrasonic surface rolling process
Materials Science & Engineering A xxx (xxxx) xxx
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Investigation of microstructural evolution in a selective laser melted Ti6Al4V alloy induced by an ultrasonic surface rolling process Zhen Wang a, Chaofeng Gao a, Zhongqiang Liu a, Zhi Wang a, Xiao Liu b, Kaho Wong c, Zhaoyao Zhou a, Zhiyu Xiao a, * a b c
National Engineering Research Center of Near-Net Shape Forming for Metallic Materials, South China University of Technology, Guangzhou, 510640, China Faculty of Science, Henan Institute of Technology, Xinxiang, 453003, China Guangzhou Zoltrix HIP Material Limited, Guangzhou, 511470, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Ti6Al4V Selective laser melting Ultrasonic surface rolling Texture evolution Dislocations
Surface modification treatment was carried out on a selective laser melted Ti6Al4V alloy using an ultrasonic surface rolling process (USRP). The results showed that a deformed layer with a gradient microstructure was produced near the surface, in which the microstructure changed from a coarse lamellar α structure to ultrafine lamellar grains, ultrafine equiaxed grains and nanograins. Moreover, there was a decrease in the intensity of the texture from the matrix to the surface. Furthermore, the dislocation density in the deformed layer first increased and then decreased from the matrix to the surface. The deformed layer exhibited a much higher hardness than the matrix, which was ascribed to grain boundary strengthening and dislocation strengthening.
1. Introduction Metallic materials often suffer from corrosion [1], wear [2] and fa tigue damage [3,4] in complex service environments, resulting in en gineering accidents and property loss. Most failures are generated at the component surface due to existing surface defects and an insufficient surface strength [5,6]. Therefore, it is important to improve the integrity of materials by optimizing the surface structure. Surface modification techniques include physical deposition methods, thermochemical sur face treatments, and severe plastic deformation [7]. Numerous studies show that the wear resistance, corrosion resistance and biocompatibility of materials, especially those of titanium alloys, can be improved by surface modification methods [8–11]. However, there are problems with physical deposition methods and thermochemical surface treatments, such as the coatings being prone to delaminating from the substrate during repeated loading, and the substrate twist during thermochemical surface treatments [7]. Surface gradient nanocrystallization on the surface of metallic components by severe plastic deformation (SPD) is an effective structural optimization approach [12], in which a deformed layer with nanocrystalline and/or ultrafine grains is formed and exhibits high strength and structural stability [13]. Useful preparation methods have been devised, including deep rolling [13–15], surface mechanical attrition treatment [16–19], shot peening [20,21], ultrasonic shot
peening [22,23], laser shock peening [24], ultrasonic impact treatment [25–27] and an ultrasonic surface rolling process (USRP) [28,29]. However, the evolution mechanisms in different materials are significantly different after SPD. Researchers have not yet reached a consensus on these evolutionary mechanism. Liu et al. [18] and Chen et al. [30] performed SPD on pure nickel and reported the formation of a plastic deformation region with a thickness greater than 350 μm. The deformed microstructures consisted of equiaxed nanograins, laminate nanograins, and submicron equiaxed and laminate grains. The grain subdivision was obtained by high-density dislocation motion. Zhu et al. [16] investigated pure titanium after SPD and suggested that the grain refinement was mainly accomplished by dislocation cells inside the microbands, where the dislocation cells were converted into subgrains upon additional strain. Liu et al. [31] concluded that nanocrystallization was affected not only by dislocation slip but also by the segmentation of twins and the nucleation of precipitated phases for recrystallization. Ao et al. [32] believed that phase transformation occurred during SPD, and grains were refined at the location of the phase transformation. In addition, Su et al. [33] revealed that vacancy-type defects also play a key role in microstructural evolution during SPD because the interaction of vacancy-type defects and dislocations causes a variation in the defect types and induces work-hardening. Therefore, the refinement mecha nisms of metallic materials are diverse, and different kinds of materials
* Corresponding author. E-mail address: [email protected] (Z. Xiao). https://doi.org/10.1016/j.msea.2019.138696 Received 10 July 2019; Received in revised form 13 November 2019; Accepted 15 November 2019 Available online 18 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: Zhen Wang, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138696
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still need to be investigated. Ti6Al4V alloys prepared by selective laser melting (SLM) are attractive materials that are very different from titanium alloys fabri cated by traditional casting or powder metallurgy methods because the SLM technique uses a laser as a heat source to selectively melt successive layers of titanium alloy powders along a set path. The microstructure of selective laser melted Ti6Al4V alloys usually consists of a number of lamellar structures with an obvious crystal orientation [34,35]. How ever, very few investigations have focused on the SPD of selective laser melted Ti6Al4V alloys, where the texture and defects, such as pores, may be mostly eliminated. In this study, a selective laser melted Ti6Al4V alloy was processed by USRP, and the effect of severe plastic deforma tion on the microstructure and hardness was studied.
Table 1 USRP parameters. Frequency (kHz)
Static load (N)
Amplitude (μm)
Spindle speed (rpm)
Axial feed (mm/rev)
30
600/750/ 900/1000
10
45
0.08
mirror polished and etched for optical microscopy (Lecia DM 5000). For electron back-scatter diffraction (EBSD) analysis, the cross-sectional plane was further polished by ion beam milling (Leica EM RES101) with a voltage of 5.0 kV and a current of 2.0 A. EBSD scans with a step size of 0.6 μm were performed using scanning electron microscopy (NOVA NanoSEM 430, USA) equipped with an Oxford Instruments Nordlys 2S detector. An analysis of the scanned data was carried out using HKL Channel 5 evaluation software. To study the detailed characteristics of the microstructure at different depths, samples were prepared for transmission electron mi croscopy (TEM) via the following steps. (i) The surface was degreased and activated with alcohol and dilute hydrochloric acid. (ii) A Ni coating was electroplated on the treated surface using an electrolytic solution (25 mL NiSO4, 10 mL NiCl2 and 10 mL C12H25SO4Na) at a constant current of 0.1 A and a temperature of 338 K. (iii) Slices were cut with a thickness of approximately 0.5 mm from the cross-section of the elec troplated sample via electrospark wire-electrode cutting, and the sam ples were punched into discs with a 3 mm diameter at the intersection of the coating and the substrate and ground to a thickness of 100 μm with sandpaper. (iv) A final thinning via double-jet electrolytic polishing (electrolyte: 590 mL CH3OHþ350 mL C6H14O2þ60 mL HClO4) was conducted at 243 K. In addition, the focused ion beam (FIB) technique was used to obtain a TEM foil at 10 μm below the treated surface. The cross-sectional microstructure was observed with a Tecnai F20 trans mission electron microscope with an accelerating voltage of 200 kV. All nanoindentation measurements were performed using a Hysitron TI-950 nanoindenter in load-controlled mode at a nominally constant load of 10000 μN.
2. Experimental methods 2.1. Materials and the USRP method Ti6Al4V alloy samples were fabricated using commercial SLM equipment (SLM solutions, 280 HL) and spherical Ti–6Al–4V ELI powder with an average size of 38 μm. The SLM processing parameters included a laser power of 250 W, scan speed of 1150 mm/s, scan spacing of 0.1 mm and layer thickness of 0.03 mm. The layers were scanned as separate stripes using a continuous laser mode and a zigzag pattern, which was alternated by 60� between each successive layer. The relative density of the selective laser melted samples was 99%. The long axis direction of the samples was perpendicular to the building direction (Fig. 1a). All of the samples were annealed in a vacuum environment at a temperature of 1273 K for 2 h before the USRP treatment. The USRP equipment consisted of two parts: an ultrasonic wave generator and a USRP operator, as shown in Fig. 1b. The working principle of the USRP equipment was as follows: the ultrasonic wave generator supplied a si nusoidal signal, the energy transducer transformed the electrical signal into mechanical energy and generated an ultrasonic vibration, and then the ultrasonic vibration was amplified with the assistance of an ampli tude changing rod, which drove a WC ball (Ф 8 mm) to strike and roll the sample. The specific parameters for the USRP are presented in Table 1.
3. Results and discussion
2.2. Microstructural characterization
3.1. Optical microscopy observation
To observe the microstructures in the treated surface, cross-sectional planes pertaining to the normal direction (ND) and rolling direction (RD) were cut from the samples. The cross-sectional planes were ground,
Fig. 2 shows the cross-sectional microstructure of the USRP-treated samples under different loads. A deformed layer was formed owing to the severe plastic deformation induced by the ultrasonic vibration. The
Fig. 1. (a) Ti6Al4V sample building direction and (b) schematic illustration of the USRP. 2
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Fig. 2. Microstructure of the USRP-treated samples under different loads: (a) 600 N, (b) 750 N, (c) 900 N and (d) 1000 N.
thickness of the layer substantially increased from ~150 μm at a load of 600 N to ~300 μm at a load of 1000 N. The matrix showed a dual microstructure composed of α/β Widmanst€ atten structures, in which the β phase was located along the boundaries of the lamellar α-phase structure. The lamellar α structure had a high aspect ratio. The lamellar α structure was curved and elongated along the rolling direction near the matrix, whereas this structure decreased in size and broke near the surface. An apparent gradient microstructure appeared on the surface of the selective laser melted Ti6Al4V alloy after the USRP treatment.
and formation of subgrain boundaries caused by the deformation [20]. However, in the case of the deformed layer near the surface (Q3), the fraction of the LAGBs decreased to ~34.2%, whereas the fraction of HAGBs increased to 62.3%, and the misorientation angle distributions were continuous. This scenario is similar to that in a previous report where the fraction of HAGBs dominated the misorientation angle dis tribution [37]. It has commonly been reported that the formation of HAGBs contributes substantially to dynamic recrystallization (DRX) [38]. However, since the temperature was far lower than that of the DRX during the USRP, then the increased fraction of HAGBs in region Q3 can only result from a large number of subgrain boundaries evolving into HAGBs under a larger deformation. Therefore, the evolution of the misorientation from the matrix to the deformed layer indicates that a significant dislocation multiplication and transition of LAGBs into HAGBs occurred as the depth from the surface decreased (see Table.2). Fig. 5 shows the Schmid factor distribution in different deformation regions, illustrating that the deformed layer had additional grains with an increased Schmid factor. The increased Schmid factor of the deformed layer indicates that more deformable grains existed in this layer than in the other regions, which was mostly owing to the decreased grain size and texture in the deformed layer after the USRP treatment.
3.2. EBSD investigation Fig. 3 shows the microstructure of the sample prepared with the USRP at a load of 1000 N, for which a gradient structure was clearly observed (Fig. 3a). To facilitate discussion, the deformed layer was marked as Q1 and Q2, and the matrix was marked as Q3 in Fig. 3a. According to the phase map (Fig. 3b), the lamellar structure comprised the hexagonal close-packed α phase and had a volume fraction of ~97%. The inverse pole figure (IPF) map (Fig. 3c) of the matrix at Q1 shows that the thickness and aspect ratio of the lamellar α structure were 9.64 μm and 2.80, respectively (summarized in Fig. 3f). The volume fraction of the lamellar α structure (mean thickness greater than 6 μm) was approximately 44%. The pole figure of Q1 shows that the lamellar α structure had a strong basal texture with a Miller index of {0001} <1010> [36], as shown in Fig. 3i. The deformed layer (the region near the matrix marked as Q2) shows that the lamellar α structures were bent and broken (Fig. 3d), and the average thickness and aspect ratio of the lamellar α structure were 5.84 μm and 2.48, respectively (Fig. 3g). The intensity of the texture was significantly decreased in the deformed layer. Furthermore, the deformed layer near the surface (Q3) shows that the structure was substantially refined and the preferred orientations of the texture shifted towards the rolling direction. The lamellar α structure was broken into a nanostructure with a thickness of 3.33 μm and an aspect ratio of 1.79 (Fig. 3h). Fig. 4 shows the misorientation distributions of the deformed layer and matrix (Q1, Q2, and Q3). Here, low-angle grain boundaries (LAGBs, <10� ), high-angle grain boundaries (HAGBs, >15� ), and medium-angle grain boundaries (MAGBs, >10� and <15� ) are depicted as red, black and green lines, respectively. As shown in Fig. 4a, the matrix (Q1) had a high fraction of LAGBs (~49.7%) and HAGBs (~44.6%). The misori entation angle distributions were discontinuous, which located at 2–20� , ~60� , and ~90� . For the deformed layer near the matrix (Q2), the fraction of the LAGBs and MAGBs increased to ~54.2% and ~19.2%, respectively, relative to that of matrix (Q1), whereas the fraction of HAGBs decreased to 26.5%. The fraction of LAGBs and MAGBs that increased at region Q2 may be related to the dislocation multiplication
3.3. TEM characterization A detailed TEM observation was performed for the USRP-treated sample under a load of 1000 N, where the TEM specimens were cho sen from the deformed layer at different depths from the surface. Fig. 6 shows TEM micrographs of the region located at a depth of 300 μm below the treated surface, which was adjacent to the strain-free matrix (Fig. 4b). A high density of dislocations was formed along the boundaries of the lamellar α structures owing to the large plastic strain induced by the USRP treatment. The corresponding selected area elec tron diffraction (SAED) pattern confirmed the hcp structure and revealed that no misorientation existed in the dislocation bands. The microstructure at this depth represented the basic features near the matrix. The microstructural characteristics at depths of 250 μm and 200 μm below the treated surface are presented in Fig. 7 (corresponding to the upper region in Fig. 4b); the dislocation density at 250 μm and 200 μm was greater than that at 300 μm. The bright-field (Fig. 7a, c) and darkfield images (Fig. 7b, d) show that ultrafine lamellar grains evolved from dislocation bands. The thickness of these ultrafine lamellar grains was a few hundred nanometers. In addition, subgrain boundaries with a misorientation of ~6� formed inside dislocation bands (Fig. 7e), which were generated by a rearrangement of edge dislocations (Fig. 7c). Under a large strain, these subgrain boundaries were transformed into the 3
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Fig. 3. EBSD characterization: (a) division of different deformation microstructures that experienced a USRP treatment of 1000 N and (b) phase map of the matrix. IPF images from (c) Q1, (d) Q2 and (e) Q3. Grain size statistics for (f) Q1, (g) Q2 and (h) Q3. Pole figures from (i) Q1, (j) Q2 and (k) Q3.
HAGBs and completed the lamellar α division. When the depth decreased to 150–100 μm (Fig. 8), ultrafine lamellar grains were divided into ultrafine equiaxed grains by transverse dislo cation walls and dislocation cells. Subsequently, the thickness of ultra fine lamellar grains was further reduced. The semicontinuous rings in the SAED pattern (inset in Fig. 8b) were obtained from circle “A”, indicating that the ultrafine equiaxed grains had random crystallo graphic orientations. Therefore, the main characteristic of the micro structure at this depth was a mixture of ultrafine lamellar grains and ultrafine equiaxed grains. Fig. 9 shows TEM micrographs for depths of 60–10 μm (corre sponding to the upper region in Fig. 4c). As observed in Fig. 9a and b, the ultrafine lamellar grains completely disappeared, and the microstruc tures were characterized by ultrafine equiaxed grains. The average thickness of the ultrafine equiaxed grains at these depths was 100–150 nm. High-density dislocation tangles existed inside ultrafine
equiaxed grains, which indicates that a high residual strain remained in this region. With a heavier deformation at the region near the surface, an additional transition of subgrain boundaries into HAGBs occurred. Consequently, equiaxed nanograins were refined from ultrafine equi axed grains at a depth of 10 μm, as shown in Fig. 9c and d, which had an average grain size of ~45 nm. A portion of the equiaxed nanograins were refined to only 5 nm–10 nm, where the lattice showed no elastic distortion (Fig. 9c inset image). 3.4. Nanoindentation of the deformation layer Fig. 10 displays the nanohardness and displacement-load curves as a function of depth from the surface after the USRP treatment under a load of 1000 N. It is interesting to note that the hardness exhibited a me chanical gradient evolution in the deformed layer, which was in accordance with the gradient structure. The deformed layer exhibited a 4
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Fig. 4. EBSD grain boundary map and misorientation angle distributions: (a, a1) Q1, (b, b1) Q2 and (c, c1) Q3. The circles represent the depth below the treated surface.
(i) Grain boundary strengthening
Table 2 Detailed statistics about the fraction of the LAGBs, MAGBs and HAGBs and the average misorientation in the three regions. Region
LAGBs, %
MAGBs, %
HAGBs, %
Average misorientation,�
Q1 Q2 Q3
49.7 54.2 34.2
5.7 19.2 3.4
44.6 26.5 62.3
25.8 18.4 46.4
The grain boundaries in the titanium alloys served as strong barriers for dislocation transmission because their hexagonal crystal structure had a limited number of independent slip modes. Hence, the titanium alloy exhibited very strong grain boundary strengthening [39]. For samples treated by the USRP, the thickness of the α-phase lamellar structure (transverse axis size) was significantly reduced from several microns near the matrix to the nanoscale near the surface. The increase in the hardness can be described by the Hall-Petch relationship [40]:
much higher hardness than the matrix, which had a hardness of ~5.57 GPa near the surface and decreased to ~4.2 GPa near the matrix, indicating that strong strain hardening occurred in the deformed layer. Fig. 10b shows the displacement-load curves, which had an elevated slope in locations near the surface, further confirming that an elevated strain hardening occurred in locations near the surface. The hardness improvement and elevated strain hardening can be explained by the following factors: grain boundary strengthening and dislocation strengthening, which can be estimated by equations for the lamellar α structure thickness and dislocation density, respectively.
H ¼ Hm þ kH dα 1=2
(1)
where H is the hardness; Hm is the matrix hardness; kH is the modified Hall-Petch constant (10.2 for titanium alloys) [40]; and dα is the average thickness of lamellar α structure. Fig. 10a shows that the measured average thickness of the lamellar structures was 45 nm at a depth of 10 nm and 390 nm at a depth of 200 μm. The corresponding hardness increases calculated by the Hall-Petch equation were ~1.51 GPa and
Fig. 5. Schmid factor for the USRP-treated surface in different regions: (a) Q1, (b) Q2 and (c) Q3. 5
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Fig. 6. (a) TEM image at a depth of 300 μm below the treated surface, (b) magnification of frame “1” and (c) corresponding SAED pattern of location “A”.
Fig. 7. (a) Bright-field (BF) image at a depth of 250 μm. (b) Magnified dark-field (DF) image of frame “1”. (c) BF and (d) DF image at a depth of 200 μm. (e) DF image corresponding to frame “2” in (a) and an inset image showing the corresponding SAED pattern of the circle “A”. (f) Inverse fast Fourier-filtered pattern of subgrain boundaries corresponding to the circle “A” in (e).
Fig. 8. (a) BF image and (b) corresponding DF image at a depth of 150 μm. (c) BF image and (d) corresponding DF image at a depth of 100 μm.
~0.51 GPa for the lamellar thicknesses of 45 nm and 390 nm, respectively.
where σ is the Taylor hardening contribution; M is the Taylor factor, which is typically between 2.0 and 3.2 for hcp titanium [41]; α is the prefactor term (here set to 0.5) [42]; G is the shear modulus of titanium (44 GPa) [42]; b is the Burgers vector (~0.3 nm) [42] and ρ is the dislocation density. The dislocation density was calculated using the following equation [43]:
(ii) Dislocation strengthening The dislocation strengthening can be predicted by the following equation [41]:
σ ¼ M αGbρ1=2
ρ ¼ b 1 ðφloc
(2)
6
φloc;err ÞΔx
1
(3)
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Fig. 9. (a) BF image and (b) corresponding DF image at a depth of 60 μm. (c) BF image and (d) corresponding DF image at a depth of 10 μm.
Fig. 10. (a) Variations in nanohardness and (b) load-displacement curves with respect to depth in the deformation layer.
where b is the Burgers vector; φloc is the averaged point-to-point misorientation, which can be determined with EBSD; φloc;err is the error in measurement of misorientation; and Δx is the step size of the EBSD measurement (0.6 μm). As shown in Fig. 10a, the dislocation density first increased and then decreased with decreasing depth. The highest dislocation density herein (3.2 � 1014 m-2) was observed at a depth of 200 μm, and the lowest dislocation density herein was 7.0 � 1012 m-2 at a depth of 10 μm. The calculated results revealed that the strength increases at depths of 200 μm and 10 μm were ~70 MPa and ~17.5 MPa, respectively. Because the hardness was proportional to the yield stress through the expression H ¼ 3σ [44], the increased hardness was converted to ~0.21 GPa and 0.05 GPa at the depths of 200 μm and 10 μm, respectively. Compared with that of the matrix, the total hardness increase was ~0.72 GPa at a depth of 200 μm and ~1.56 GPa at a depth of 10 μm, which was in accordance with the measured nanohardness results. These results also indicate that grain boundary strengthening had a greater effect on the hardness than dislocation strengthening.
2. The fraction of LAGBs and HAGBs presented an “up-to-down” trend with decreasing depth below the treated surface. The highest dislo cation density herein of approximately 7.0 � 1012 m-2 occurred at a depth of 200 μm below the processed surface. A decreased disloca tion density and an increased number of HAGBs were observed near the surface, which was attributed to the significant transformation of LAGBs (dislocations) into HAGBs along with the depth from the surface. 4. The dislocation motion played the most important role in the refinement of the lamellar α structure in selective laser melted Ti6Al4V alloy as a result of subgrain boundaries being formed inside the dislocation bands and transforming into HAGBs. The average size of the equiaxed grains at a depth of 10 μm below the processed surface was approximately 45 nm. 5. The deformed layer exhibited a mechanical gradient distribution where the outer region exhibited a much higher hardness than the inner region of the deformed layer and matrix, in which the hardness in the deformed layer ranged from 5.57 GPa near the surface to 4.2 GPa near the matrix. A strong strain hardening effect was observed in the deformed layer. The strengthening mechanism can be ascribed to grain boundary strengthening and dislocation strengthening.
4. Conclusions In this work, the ultrasonic surface rolling process was performed on a selective laser melted Ti6Al4V alloy, leading to a deformed layer with a gradient structure and significantly improved hardness. The main conclusions are summarized as follows:
Author contributions The author Zhen Wang did the experiment and wrote the article; the authors Chaofeng Gao and Zhongqiang Liu collected data; the author Zhi Wang improved the manuscript language; the authors Xiao Liu, Kaho Wong and Zhaoyao Zhou gave technical supports; The author Zhiyu Xiao provided helpful discussions and modification on the article.
1. A gradient structure deformed layer with a thickness of 300 μm was obtained when selective laser melted Ti6Al4V alloy was subjected to the USRP. From the matrix to the surface, the lamellar structure was broken, and both the average thickness and aspect ratio decreased. The gradient structure consisted of high-density dislocations, ultra fine lamellar grains, ultrafine equiaxed grains and equiaxed nanograins.
Declaration of competing interest The authors declare no competing financial interests. 7
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Acknowledgments
[22] G. Liu, J. Lu, K. Lu, Surface nanocrystallization of 316L stainless steel induced by ultrasonic shot peening, Mater. Sci. Eng. A 286 (2000) 91–95, https://doi.org/ 10.1016/S0921-5093(00)00686-9. [23] N.R. Tao, M.L. Sui, J. Lu, K. Lua, Surface nanocrystallization of iron induced by ultrasonic shot peening, Nanostructured Mater. 11 (1999) 433–440, https://doi. org/10.1016/S0965-9773(99)00324-4. [24] S. Lou, Y. Li, L. Zhou, X. Nie, G. He, W. He, Surface nanocrystallization of metallic alloys with different stacking fault energy induced by laser shock processing, Mater. Des. 104 (2016) 320–326, https://doi.org/10.1016/j.matdes.2016.05.028. [25] A.I. Dekhtyar, B.N. Mordyuk, D.G. Savvakin, V.I. Bondarchuk, I.V. Moiseeva, N. I. Khripta, Enhanced fatigue behavior of powder metallurgy Ti-6Al-4V alloy by applying ultrasonic impact treatment, Mater. Sci. Eng. A 641 (2015) 348–359, https://doi.org/10.1016/j.msea.2015.06.072. [26] B.N. Mordyuk, M.O. Iefimov, G.I. Prokopenko, T.V. Golub, M.I. Danylenko, Structure, microhardness and damping characteristics of Al matrix composite reinforced with AlCuFe or Ti using ultrasonic impact peening, Surf. Coat. Technol. 204 (2010) 1590–1598, https://doi.org/10.1016/j.surfcoat.2009.10.009. [27] B.N. Mordyuk, O.P. Karasevskaya, G.I. Prokopenko, Structurally induced enhancement in corrosion resistance of Zr-2.5%Nb alloy in saline solution by applying ultrasonic impact peening, Mater. Sci. Eng. A 559 (2013) 453–461, https://doi.org/10.1016/j.msea.2012.08.125. [28] Z. Wang, Z. Xiao, C. Huang, L. Wen, W. Zhang, Influence of ultrasonic surface rolling on microstructure and wear behavior of selective laser melted Ti-6Al-4V alloy, Materials 10 (2017) 1203, https://doi.org/10.3390/ma10101203. [29] G. Li, S.G. Qu, Y.X. Pan, X.Q. Li, Effects of the different frequencies and loads of ultrasonic surface rolling on surface mechanical properties and fretting wear resistance of HIP Ti–6Al–4V alloy, Appl. Surf. Sci. 389 (2016) 324–334, https:// doi.org/10.1016/j.apsusc.2016.07.120. [30] L. Chen, X. Ren, W. Zhou, Z. Tong, S. Adu-Gyamfi, Y. Ye, Y. Ren, Evolution of microstructure and grain refinement mechanism of pure nickel induced by laser shock peening, Mater. Sci. Eng. A 728 (2018) 20–29, https://doi.org/10.1016/j. msea.2018.04.105. [31] D. Liu, D. Liu, X. Zhang, C. Liu, N. Ao, Surface nanocrystallization of 17-4 precipitation-hardening stainless steel subjected to ultrasonic surface rolling process, Mater. Sci. Eng. A 726 (2018) 69–81, https://doi.org/10.1016/j. msea.2018.04.033. [32] N. Ao, D. Liu, X. Xu, X. Zhang, D. Liu, Gradient nanostructure evolution and phase transformation of α phase in Ti-6Al-4V alloy induced by ultrasonic surface rolling process, Mater. Sci. Eng. A 742 (2018) 820–834, https://doi.org/10.1016/j. msea.2018.10.098. [33] L.H. Su, C. Lu, L.Z. He, L.C. Zhang, P. Guagliardo, A.K. Tieu, S.N. Samarin, J. F. Williams, H.J. Li, Study of vacancy-type defects by positron annihilation in ultrafine-grained aluminum severely deformed at room and cryogenic temperatures, Acta Mater. 60 (2012) 4218–4228, https://doi.org/10.1016/j. actamat.2012.04.003. [34] M. Seifi, A. Salem, J. Beuth, O. Harrysson, J.J. Lewandowski, Overview of materials qualification needs for metal additive manufacturing, J. Occup. Med. (2016) 747–764, https://doi.org/10.1007/s11837-015-1810-0. [35] M. Simonelli, Y.Y. Tse, C. Tuck, On the texture formation of selective laser melted Ti-6Al-4V, Metall. Mater. Trans. A (2014), https://doi.org/10.1007/s11661-0142218-0. [36] C. Yang, J.A. Lin, Y.F. Ding, W.W. Zhang, Y.Y. Li, Z.Q. Fu, F. Chen, E.J. Lavernia, Texture evolution and mechanical behavior of commercially pure Ti processed via pulsed electric current treatment, J. Mater. Sci. 51 (2016) 10608–10619, https:// doi.org/10.1007/s10853-016-0282-0. [37] L. Wang, L. Xie, Y. Lv, L.C. Zhang, L. Chen, Q. Meng, J. Qu, D. Zhang, W. Lu, Microstructure evolution and superelastic behavior in Ti-35Nb-2Ta-3Zr alloy processed by friction stir processing, Acta Mater. 131 (2017) 499–510, https://doi. org/10.1016/j.actamat.2017.03.079. [38] Z. Liu, P. Li, L. Xiong, T. Liu, L. He, High-temperature tensile deformation behavior and microstructure evolution of Ti55 titanium alloy, Mater. Sci. Eng. A 680 (2017) 259–269, https://doi.org/10.1016/j.msea.2016.10.095. [39] Z. Ji, H. Yang, H. Li, Contributions of microstructural features to the integrated hardness of TA15 titanium alloy, Mater. Sci. Eng. A 628 (2015) 358–365, https:// doi.org/10.1016/j.msea.2014.12.087. [40] C.Y. Hyun, J.H. Lee, H.K. Kim, Microstructures and mechanical properties of ultrafine grained pure Ti produced by severe plastic deformation, Res. Chem. Intermed. 36 (2010) 629–638, https://doi.org/10.1007/s11164-010-0198-2. [41] P. Luo, D.T. McDonald, W. Xu, S. Palanisamy, M.S. Dargusch, K. Xia, A modified Hall-Petch relationship in ultrafine-grained titanium recycled from chips by equal channel angular pressing, Scr. Mater. 66 (2012) 785–788, https://doi.org/ 10.1016/j.scriptamat.2012.02.008. [42] B.J. Hayes, B.W. Martin, B. Welk, S.J. Kuhr, T.K. Ales, D.A. Brice, I. Ghamarian, A. H. Baker, C.V. Haden, D.G. Harlow, H.L. Fraser, P.C. Collins, Predicting tensile properties of Ti-6Al-4V produced via directed energy deposition, Acta Mater. 133 (2017) 120–133, https://doi.org/10.1016/j.actamat.2017.05.025. [43] R. Kapoor, N. Kumar, R.S. Mishra, C.S. Huskamp, K.K. Sankaran, Influence of fraction of high angle boundaries on the mechanical behavior of an ultrafine grained Al-Mg alloy, Mater. Sci. Eng. A 527 (2010) 5246–5254, https://doi.org/ 10.1016/j.msea.2010.04.086. [44] J. Moon, S. Kim, J. il Jang, J. Lee, C. Lee, Orowan strengthening effect on the nanoindentation hardness of the ferrite matrix in microalloyed steels, Mater. Sci. Eng. A 487 (2008) 552–557, https://doi.org/10.1016/j.msea.2007.10.046.
This work was financially supported by the National Natural Science Foundation of China (Grant No. 51627805), the Natural Science Foun dation Project of Guangdong Province (Grant No. 2015A030312003), and the Science and Technology Research Projects of Guangdong Province (Grant No. 2014B010129003, 2015B020238008, 2016B090931006, 2017B090901025). References [1] C.T. dos Santos, C. Barbosa, M.J. Monteiro, I.C. Abud, I.M.V. Caminha, C.R. M. Roesler, Characterization of the fretting corrosion behavior, surface and debris from head-taper interface of two different modular hip prostheses, J. Mech. Behav. Biomed. Mater. 62 (2016) 71–82, https://doi.org/10.1016/j.jmbbm.2016.04.036. [2] D. Gu, Y. Hagedorn, W. Meiners, G. Meng, Densification behavior , microstructure evolution , and wear performance of selective laser melting processed commercially pure titanium, Acta Mater. 60 (2012) 3849–3860, https://doi.org/ 10.1016/j.actamat.2012.04.006. [3] G. Kasperovich, J. Hausmann, Improvement of fatigue resistance and ductility of TiAl6V4 processed by selective laser melting, J. Mater. Process. Technol. 220 (2015) 202–214, https://doi.org/10.1016/j.jmatprotec.2015.01.025. [4] Y.H. Lu, Z.Y. Xiao, L. Hu, F. Luo, Y.B. Wu, D.H. Ni, Ultra-high cycle fatigue behaviour of warm compaction Fe-Cu-Ni-Mo-C sintered material, Mater. Des. 55 (2014) 758–763, https://doi.org/10.1016/j.matdes.2013.10.049. [5] I. Serrano-Munoz, J.Y. Buffiere, R. Mokso, C. Verdu, Y. Nadot, Location, location &size: defects close to surfaces dominate fatigue crack initiation, Sci. Rep. 7 (2017), 45239, https://doi.org/10.1038/srep45239. [6] R. Sonntag, J. Reinders, J. Gibmeier, J.P. Kretzer, Fatigue performance of medical Ti6Al4V alloy after mechanical surface treatments, PLoS One 10 (2015), e0121963, https://doi.org/10.1371/journal.pone.0121963. [7] L.C. Zhang, L.Y. Chen, A review on biomedical titanium alloys: recent progress and prospect, Adv. Eng. Mater. 21 (2019) 1–29, https://doi.org/10.1002/ adem.201801215. [8] B. Subramanian, C.V. Muraleedharan, R. Ananthakumar, M. Jayachandran, A comparative study of titanium nitride (TiN), titanium oxy nitride (TiON) and titanium aluminum nitride (TiAlN), as surface coatings for bio implants, Surf. Coat. Technol. 205 (2011) 5014–5020, https://doi.org/10.1016/j.surfcoat.2011.05.004. [9] S. Kumar, T.S.N.S. Narayanan, S. Ganesh Sundara Raman, S.K. Seshadri, Surface modification of CP-Ti to improve the fretting-corrosion resistance: thermal oxidation vs. anodizing, Mater. Sci. Eng. C 30 (2010) 921–927, https://doi.org/ 10.1016/j.msec.2010.03.024. [10] A. Bloyce, P.-Y. Qi, H. Dong, T. Bell, Surface modification of titanium alloys for combined improvements in corrosion and wear resistance, Surf. Coat. Technol. 107 (1998) 125–132, https://doi.org/10.1016/S0257-8972(98)00580-5. [11] M. Niinomi, Recent metallic materials for biomedical applications, Metall. Mater. Trans. A 33 (2002) 477–486, https://doi.org/10.1007/s11661-002-0109-2. [12] X.H. Chen, J. Lu, L. Lu, K. Lu, Tensile properties of a nanocrystalline 316L austenitic stainless steel, Scr. Mater. 52 (2005) 1039–1044, https://doi.org/ 10.1016/J.SCRIPTAMAT.2005.01.023. [13] H. Gleiter, Nanostructured materials: basic concepts and microstructure, Acta Mater. 48 (2000) 1–29, https://doi.org/10.1016/S1359-6454(99)00285-2. [14] R.K. Nalla, I. Altenberger, U. Noster, G.Y. Liu, B. Scholtes, R.O. Ritchie, On the influence of mechanical surface treatments-deep rolling and laser shock peening-on the fatigue behavior of Ti-6Al-4V at ambient and elevated temperatures, Mater. Sci. Eng. A 355 (2003) 216–230, https://doi.org/10.1016/S0921-5093(03)000698. [15] X. Liu, Z.y. Xiao, H.j. Guan, W. Zhang, F.l. Li, Friction and wear behaviours of surface densified powder metallurgy Fe–2Cu–0.6C material, Powder Metall. 59 (2016) 329–334, https://doi.org/10.1080/00325899.2016.1242880. [16] K.Y. Zhu, A. Vassel, F. Brisset, K. Lu, J. Lu, Nanostructure formation mechanism of α-titanium using SMAT, Acta Mater. 52 (2004) 4101–4110, https://doi.org/ 10.1016/j.actamat.2004.05.023. [17] L. Jin, W. Cui, X. Song, L. Zhou, The formation mechanisms of surface nanocrystallites in β-type biomedical TiNbZrFe alloy by surface mechanical attrition treatment, Appl. Surf. Sci. 347 (2015) 553–560, https://doi.org/10.1016/ j.apsusc.2015.04.137. [18] X.C. Liu, H.W. Zhang, K. Lu, Formation of nano-laminated structure in nickel by means of surface mechanical grinding treatment, Acta Mater. 96 (2015) 24–36, https://doi.org/10.1016/j.actamat.2015.06.014. [19] W.L. Li, N.R. Tao, K. Lu, Fabrication of a gradient nano-micro-structured surface layer on bulk copper by means of a surface mechanical grinding treatment, Scr. Mater. 59 (2008) 546–549, https://doi.org/10.1016/j.scriptamat.2008.05.003. [20] Y.G. Liu, M.Q. Li, H.J. Liu, Surface nanocrystallization and gradient structure developed in the bulk TC4 alloy processed by shot peening, J. Alloy. Comp. 685 (2016) 186–193, https://doi.org/10.1016/j.jallcom.2016.05.295. [21] H. Li, Y. Liu, M. Li, H. Liu, The gradient crystalline structure and microhardness in the treated layer of TC17 via high energy shot peening, Appl. Surf. Sci. 357 (2015) 197–203, https://doi.org/10.1016/j.apsusc.2015.09.013.
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Investigation of microstructural evolution in a selective laser melted Ti6Al4V alloy induced by an ultrasonic surface rolling process Zhen Wang a, Chaofeng Gao a, Zhongqiang Liu a, Zhi Wang a, Xiao Liu b, Kaho Wong c, Zhaoyao Zhou a, Zhiyu Xiao a, * a b c
National Engineering Research Center of Near-Net Shape Forming for Metallic Materials, South China University of Technology, Guangzhou, 510640, China Faculty of Science, Henan Institute of Technology, Xinxiang, 453003, China Guangzhou Zoltrix HIP Material Limited, Guangzhou, 511470, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Ti6Al4V Selective laser melting Ultrasonic surface rolling Texture evolution Dislocations
Surface modification treatment was carried out on a selective laser melted Ti6Al4V alloy using an ultrasonic surface rolling process (USRP). The results showed that a deformed layer with a gradient microstructure was produced near the surface, in which the microstructure changed from a coarse lamellar α structure to ultrafine lamellar grains, ultrafine equiaxed grains and nanograins. Moreover, there was a decrease in the intensity of the texture from the matrix to the surface. Furthermore, the dislocation density in the deformed layer first increased and then decreased from the matrix to the surface. The deformed layer exhibited a much higher hardness than the matrix, which was ascribed to grain boundary strengthening and dislocation strengthening.
1. Introduction Metallic materials often suffer from corrosion [1], wear [2] and fa tigue damage [3,4] in complex service environments, resulting in en gineering accidents and property loss. Most failures are generated at the component surface due to existing surface defects and an insufficient surface strength [5,6]. Therefore, it is important to improve the integrity of materials by optimizing the surface structure. Surface modification techniques include physical deposition methods, thermochemical sur face treatments, and severe plastic deformation [7]. Numerous studies show that the wear resistance, corrosion resistance and biocompatibility of materials, especially those of titanium alloys, can be improved by surface modification methods [8–11]. However, there are problems with physical deposition methods and thermochemical surface treatments, such as the coatings being prone to delaminating from the substrate during repeated loading, and the substrate twist during thermochemical surface treatments [7]. Surface gradient nanocrystallization on the surface of metallic components by severe plastic deformation (SPD) is an effective structural optimization approach [12], in which a deformed layer with nanocrystalline and/or ultrafine grains is formed and exhibits high strength and structural stability [13]. Useful preparation methods have been devised, including deep rolling [13–15], surface mechanical attrition treatment [16–19], shot peening [20,21], ultrasonic shot
peening [22,23], laser shock peening [24], ultrasonic impact treatment [25–27] and an ultrasonic surface rolling process (USRP) [28,29]. However, the evolution mechanisms in different materials are significantly different after SPD. Researchers have not yet reached a consensus on these evolutionary mechanism. Liu et al. [18] and Chen et al. [30] performed SPD on pure nickel and reported the formation of a plastic deformation region with a thickness greater than 350 μm. The deformed microstructures consisted of equiaxed nanograins, laminate nanograins, and submicron equiaxed and laminate grains. The grain subdivision was obtained by high-density dislocation motion. Zhu et al. [16] investigated pure titanium after SPD and suggested that the grain refinement was mainly accomplished by dislocation cells inside the microbands, where the dislocation cells were converted into subgrains upon additional strain. Liu et al. [31] concluded that nanocrystallization was affected not only by dislocation slip but also by the segmentation of twins and the nucleation of precipitated phases for recrystallization. Ao et al. [32] believed that phase transformation occurred during SPD, and grains were refined at the location of the phase transformation. In addition, Su et al. [33] revealed that vacancy-type defects also play a key role in microstructural evolution during SPD because the interaction of vacancy-type defects and dislocations causes a variation in the defect types and induces work-hardening. Therefore, the refinement mecha nisms of metallic materials are diverse, and different kinds of materials
* Corresponding author. E-mail address: [email protected] (Z. Xiao). https://doi.org/10.1016/j.msea.2019.138696 Received 10 July 2019; Received in revised form 13 November 2019; Accepted 15 November 2019 Available online 18 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: Zhen Wang, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138696
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still need to be investigated. Ti6Al4V alloys prepared by selective laser melting (SLM) are attractive materials that are very different from titanium alloys fabri cated by traditional casting or powder metallurgy methods because the SLM technique uses a laser as a heat source to selectively melt successive layers of titanium alloy powders along a set path. The microstructure of selective laser melted Ti6Al4V alloys usually consists of a number of lamellar structures with an obvious crystal orientation [34,35]. How ever, very few investigations have focused on the SPD of selective laser melted Ti6Al4V alloys, where the texture and defects, such as pores, may be mostly eliminated. In this study, a selective laser melted Ti6Al4V alloy was processed by USRP, and the effect of severe plastic deforma tion on the microstructure and hardness was studied.
Table 1 USRP parameters. Frequency (kHz)
Static load (N)
Amplitude (μm)
Spindle speed (rpm)
Axial feed (mm/rev)
30
600/750/ 900/1000
10
45
0.08
mirror polished and etched for optical microscopy (Lecia DM 5000). For electron back-scatter diffraction (EBSD) analysis, the cross-sectional plane was further polished by ion beam milling (Leica EM RES101) with a voltage of 5.0 kV and a current of 2.0 A. EBSD scans with a step size of 0.6 μm were performed using scanning electron microscopy (NOVA NanoSEM 430, USA) equipped with an Oxford Instruments Nordlys 2S detector. An analysis of the scanned data was carried out using HKL Channel 5 evaluation software. To study the detailed characteristics of the microstructure at different depths, samples were prepared for transmission electron mi croscopy (TEM) via the following steps. (i) The surface was degreased and activated with alcohol and dilute hydrochloric acid. (ii) A Ni coating was electroplated on the treated surface using an electrolytic solution (25 mL NiSO4, 10 mL NiCl2 and 10 mL C12H25SO4Na) at a constant current of 0.1 A and a temperature of 338 K. (iii) Slices were cut with a thickness of approximately 0.5 mm from the cross-section of the elec troplated sample via electrospark wire-electrode cutting, and the sam ples were punched into discs with a 3 mm diameter at the intersection of the coating and the substrate and ground to a thickness of 100 μm with sandpaper. (iv) A final thinning via double-jet electrolytic polishing (electrolyte: 590 mL CH3OHþ350 mL C6H14O2þ60 mL HClO4) was conducted at 243 K. In addition, the focused ion beam (FIB) technique was used to obtain a TEM foil at 10 μm below the treated surface. The cross-sectional microstructure was observed with a Tecnai F20 trans mission electron microscope with an accelerating voltage of 200 kV. All nanoindentation measurements were performed using a Hysitron TI-950 nanoindenter in load-controlled mode at a nominally constant load of 10000 μN.
2. Experimental methods 2.1. Materials and the USRP method Ti6Al4V alloy samples were fabricated using commercial SLM equipment (SLM solutions, 280 HL) and spherical Ti–6Al–4V ELI powder with an average size of 38 μm. The SLM processing parameters included a laser power of 250 W, scan speed of 1150 mm/s, scan spacing of 0.1 mm and layer thickness of 0.03 mm. The layers were scanned as separate stripes using a continuous laser mode and a zigzag pattern, which was alternated by 60� between each successive layer. The relative density of the selective laser melted samples was 99%. The long axis direction of the samples was perpendicular to the building direction (Fig. 1a). All of the samples were annealed in a vacuum environment at a temperature of 1273 K for 2 h before the USRP treatment. The USRP equipment consisted of two parts: an ultrasonic wave generator and a USRP operator, as shown in Fig. 1b. The working principle of the USRP equipment was as follows: the ultrasonic wave generator supplied a si nusoidal signal, the energy transducer transformed the electrical signal into mechanical energy and generated an ultrasonic vibration, and then the ultrasonic vibration was amplified with the assistance of an ampli tude changing rod, which drove a WC ball (Ф 8 mm) to strike and roll the sample. The specific parameters for the USRP are presented in Table 1.
3. Results and discussion
2.2. Microstructural characterization
3.1. Optical microscopy observation
To observe the microstructures in the treated surface, cross-sectional planes pertaining to the normal direction (ND) and rolling direction (RD) were cut from the samples. The cross-sectional planes were ground,
Fig. 2 shows the cross-sectional microstructure of the USRP-treated samples under different loads. A deformed layer was formed owing to the severe plastic deformation induced by the ultrasonic vibration. The
Fig. 1. (a) Ti6Al4V sample building direction and (b) schematic illustration of the USRP. 2
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Fig. 2. Microstructure of the USRP-treated samples under different loads: (a) 600 N, (b) 750 N, (c) 900 N and (d) 1000 N.
thickness of the layer substantially increased from ~150 μm at a load of 600 N to ~300 μm at a load of 1000 N. The matrix showed a dual microstructure composed of α/β Widmanst€ atten structures, in which the β phase was located along the boundaries of the lamellar α-phase structure. The lamellar α structure had a high aspect ratio. The lamellar α structure was curved and elongated along the rolling direction near the matrix, whereas this structure decreased in size and broke near the surface. An apparent gradient microstructure appeared on the surface of the selective laser melted Ti6Al4V alloy after the USRP treatment.
and formation of subgrain boundaries caused by the deformation [20]. However, in the case of the deformed layer near the surface (Q3), the fraction of the LAGBs decreased to ~34.2%, whereas the fraction of HAGBs increased to 62.3%, and the misorientation angle distributions were continuous. This scenario is similar to that in a previous report where the fraction of HAGBs dominated the misorientation angle dis tribution [37]. It has commonly been reported that the formation of HAGBs contributes substantially to dynamic recrystallization (DRX) [38]. However, since the temperature was far lower than that of the DRX during the USRP, then the increased fraction of HAGBs in region Q3 can only result from a large number of subgrain boundaries evolving into HAGBs under a larger deformation. Therefore, the evolution of the misorientation from the matrix to the deformed layer indicates that a significant dislocation multiplication and transition of LAGBs into HAGBs occurred as the depth from the surface decreased (see Table.2). Fig. 5 shows the Schmid factor distribution in different deformation regions, illustrating that the deformed layer had additional grains with an increased Schmid factor. The increased Schmid factor of the deformed layer indicates that more deformable grains existed in this layer than in the other regions, which was mostly owing to the decreased grain size and texture in the deformed layer after the USRP treatment.
3.2. EBSD investigation Fig. 3 shows the microstructure of the sample prepared with the USRP at a load of 1000 N, for which a gradient structure was clearly observed (Fig. 3a). To facilitate discussion, the deformed layer was marked as Q1 and Q2, and the matrix was marked as Q3 in Fig. 3a. According to the phase map (Fig. 3b), the lamellar structure comprised the hexagonal close-packed α phase and had a volume fraction of ~97%. The inverse pole figure (IPF) map (Fig. 3c) of the matrix at Q1 shows that the thickness and aspect ratio of the lamellar α structure were 9.64 μm and 2.80, respectively (summarized in Fig. 3f). The volume fraction of the lamellar α structure (mean thickness greater than 6 μm) was approximately 44%. The pole figure of Q1 shows that the lamellar α structure had a strong basal texture with a Miller index of {0001} <1010> [36], as shown in Fig. 3i. The deformed layer (the region near the matrix marked as Q2) shows that the lamellar α structures were bent and broken (Fig. 3d), and the average thickness and aspect ratio of the lamellar α structure were 5.84 μm and 2.48, respectively (Fig. 3g). The intensity of the texture was significantly decreased in the deformed layer. Furthermore, the deformed layer near the surface (Q3) shows that the structure was substantially refined and the preferred orientations of the texture shifted towards the rolling direction. The lamellar α structure was broken into a nanostructure with a thickness of 3.33 μm and an aspect ratio of 1.79 (Fig. 3h). Fig. 4 shows the misorientation distributions of the deformed layer and matrix (Q1, Q2, and Q3). Here, low-angle grain boundaries (LAGBs, <10� ), high-angle grain boundaries (HAGBs, >15� ), and medium-angle grain boundaries (MAGBs, >10� and <15� ) are depicted as red, black and green lines, respectively. As shown in Fig. 4a, the matrix (Q1) had a high fraction of LAGBs (~49.7%) and HAGBs (~44.6%). The misori entation angle distributions were discontinuous, which located at 2–20� , ~60� , and ~90� . For the deformed layer near the matrix (Q2), the fraction of the LAGBs and MAGBs increased to ~54.2% and ~19.2%, respectively, relative to that of matrix (Q1), whereas the fraction of HAGBs decreased to 26.5%. The fraction of LAGBs and MAGBs that increased at region Q2 may be related to the dislocation multiplication
3.3. TEM characterization A detailed TEM observation was performed for the USRP-treated sample under a load of 1000 N, where the TEM specimens were cho sen from the deformed layer at different depths from the surface. Fig. 6 shows TEM micrographs of the region located at a depth of 300 μm below the treated surface, which was adjacent to the strain-free matrix (Fig. 4b). A high density of dislocations was formed along the boundaries of the lamellar α structures owing to the large plastic strain induced by the USRP treatment. The corresponding selected area elec tron diffraction (SAED) pattern confirmed the hcp structure and revealed that no misorientation existed in the dislocation bands. The microstructure at this depth represented the basic features near the matrix. The microstructural characteristics at depths of 250 μm and 200 μm below the treated surface are presented in Fig. 7 (corresponding to the upper region in Fig. 4b); the dislocation density at 250 μm and 200 μm was greater than that at 300 μm. The bright-field (Fig. 7a, c) and darkfield images (Fig. 7b, d) show that ultrafine lamellar grains evolved from dislocation bands. The thickness of these ultrafine lamellar grains was a few hundred nanometers. In addition, subgrain boundaries with a misorientation of ~6� formed inside dislocation bands (Fig. 7e), which were generated by a rearrangement of edge dislocations (Fig. 7c). Under a large strain, these subgrain boundaries were transformed into the 3
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Fig. 3. EBSD characterization: (a) division of different deformation microstructures that experienced a USRP treatment of 1000 N and (b) phase map of the matrix. IPF images from (c) Q1, (d) Q2 and (e) Q3. Grain size statistics for (f) Q1, (g) Q2 and (h) Q3. Pole figures from (i) Q1, (j) Q2 and (k) Q3.
HAGBs and completed the lamellar α division. When the depth decreased to 150–100 μm (Fig. 8), ultrafine lamellar grains were divided into ultrafine equiaxed grains by transverse dislo cation walls and dislocation cells. Subsequently, the thickness of ultra fine lamellar grains was further reduced. The semicontinuous rings in the SAED pattern (inset in Fig. 8b) were obtained from circle “A”, indicating that the ultrafine equiaxed grains had random crystallo graphic orientations. Therefore, the main characteristic of the micro structure at this depth was a mixture of ultrafine lamellar grains and ultrafine equiaxed grains. Fig. 9 shows TEM micrographs for depths of 60–10 μm (corre sponding to the upper region in Fig. 4c). As observed in Fig. 9a and b, the ultrafine lamellar grains completely disappeared, and the microstruc tures were characterized by ultrafine equiaxed grains. The average thickness of the ultrafine equiaxed grains at these depths was 100–150 nm. High-density dislocation tangles existed inside ultrafine
equiaxed grains, which indicates that a high residual strain remained in this region. With a heavier deformation at the region near the surface, an additional transition of subgrain boundaries into HAGBs occurred. Consequently, equiaxed nanograins were refined from ultrafine equi axed grains at a depth of 10 μm, as shown in Fig. 9c and d, which had an average grain size of ~45 nm. A portion of the equiaxed nanograins were refined to only 5 nm–10 nm, where the lattice showed no elastic distortion (Fig. 9c inset image). 3.4. Nanoindentation of the deformation layer Fig. 10 displays the nanohardness and displacement-load curves as a function of depth from the surface after the USRP treatment under a load of 1000 N. It is interesting to note that the hardness exhibited a me chanical gradient evolution in the deformed layer, which was in accordance with the gradient structure. The deformed layer exhibited a 4
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Fig. 4. EBSD grain boundary map and misorientation angle distributions: (a, a1) Q1, (b, b1) Q2 and (c, c1) Q3. The circles represent the depth below the treated surface.
(i) Grain boundary strengthening
Table 2 Detailed statistics about the fraction of the LAGBs, MAGBs and HAGBs and the average misorientation in the three regions. Region
LAGBs, %
MAGBs, %
HAGBs, %
Average misorientation,�
Q1 Q2 Q3
49.7 54.2 34.2
5.7 19.2 3.4
44.6 26.5 62.3
25.8 18.4 46.4
The grain boundaries in the titanium alloys served as strong barriers for dislocation transmission because their hexagonal crystal structure had a limited number of independent slip modes. Hence, the titanium alloy exhibited very strong grain boundary strengthening [39]. For samples treated by the USRP, the thickness of the α-phase lamellar structure (transverse axis size) was significantly reduced from several microns near the matrix to the nanoscale near the surface. The increase in the hardness can be described by the Hall-Petch relationship [40]:
much higher hardness than the matrix, which had a hardness of ~5.57 GPa near the surface and decreased to ~4.2 GPa near the matrix, indicating that strong strain hardening occurred in the deformed layer. Fig. 10b shows the displacement-load curves, which had an elevated slope in locations near the surface, further confirming that an elevated strain hardening occurred in locations near the surface. The hardness improvement and elevated strain hardening can be explained by the following factors: grain boundary strengthening and dislocation strengthening, which can be estimated by equations for the lamellar α structure thickness and dislocation density, respectively.
H ¼ Hm þ kH dα 1=2
(1)
where H is the hardness; Hm is the matrix hardness; kH is the modified Hall-Petch constant (10.2 for titanium alloys) [40]; and dα is the average thickness of lamellar α structure. Fig. 10a shows that the measured average thickness of the lamellar structures was 45 nm at a depth of 10 nm and 390 nm at a depth of 200 μm. The corresponding hardness increases calculated by the Hall-Petch equation were ~1.51 GPa and
Fig. 5. Schmid factor for the USRP-treated surface in different regions: (a) Q1, (b) Q2 and (c) Q3. 5
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Fig. 6. (a) TEM image at a depth of 300 μm below the treated surface, (b) magnification of frame “1” and (c) corresponding SAED pattern of location “A”.
Fig. 7. (a) Bright-field (BF) image at a depth of 250 μm. (b) Magnified dark-field (DF) image of frame “1”. (c) BF and (d) DF image at a depth of 200 μm. (e) DF image corresponding to frame “2” in (a) and an inset image showing the corresponding SAED pattern of the circle “A”. (f) Inverse fast Fourier-filtered pattern of subgrain boundaries corresponding to the circle “A” in (e).
Fig. 8. (a) BF image and (b) corresponding DF image at a depth of 150 μm. (c) BF image and (d) corresponding DF image at a depth of 100 μm.
~0.51 GPa for the lamellar thicknesses of 45 nm and 390 nm, respectively.
where σ is the Taylor hardening contribution; M is the Taylor factor, which is typically between 2.0 and 3.2 for hcp titanium [41]; α is the prefactor term (here set to 0.5) [42]; G is the shear modulus of titanium (44 GPa) [42]; b is the Burgers vector (~0.3 nm) [42] and ρ is the dislocation density. The dislocation density was calculated using the following equation [43]:
(ii) Dislocation strengthening The dislocation strengthening can be predicted by the following equation [41]:
σ ¼ M αGbρ1=2
ρ ¼ b 1 ðφloc
(2)
6
φloc;err ÞΔx
1
(3)
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Fig. 9. (a) BF image and (b) corresponding DF image at a depth of 60 μm. (c) BF image and (d) corresponding DF image at a depth of 10 μm.
Fig. 10. (a) Variations in nanohardness and (b) load-displacement curves with respect to depth in the deformation layer.
where b is the Burgers vector; φloc is the averaged point-to-point misorientation, which can be determined with EBSD; φloc;err is the error in measurement of misorientation; and Δx is the step size of the EBSD measurement (0.6 μm). As shown in Fig. 10a, the dislocation density first increased and then decreased with decreasing depth. The highest dislocation density herein (3.2 � 1014 m-2) was observed at a depth of 200 μm, and the lowest dislocation density herein was 7.0 � 1012 m-2 at a depth of 10 μm. The calculated results revealed that the strength increases at depths of 200 μm and 10 μm were ~70 MPa and ~17.5 MPa, respectively. Because the hardness was proportional to the yield stress through the expression H ¼ 3σ [44], the increased hardness was converted to ~0.21 GPa and 0.05 GPa at the depths of 200 μm and 10 μm, respectively. Compared with that of the matrix, the total hardness increase was ~0.72 GPa at a depth of 200 μm and ~1.56 GPa at a depth of 10 μm, which was in accordance with the measured nanohardness results. These results also indicate that grain boundary strengthening had a greater effect on the hardness than dislocation strengthening.
2. The fraction of LAGBs and HAGBs presented an “up-to-down” trend with decreasing depth below the treated surface. The highest dislo cation density herein of approximately 7.0 � 1012 m-2 occurred at a depth of 200 μm below the processed surface. A decreased disloca tion density and an increased number of HAGBs were observed near the surface, which was attributed to the significant transformation of LAGBs (dislocations) into HAGBs along with the depth from the surface. 4. The dislocation motion played the most important role in the refinement of the lamellar α structure in selective laser melted Ti6Al4V alloy as a result of subgrain boundaries being formed inside the dislocation bands and transforming into HAGBs. The average size of the equiaxed grains at a depth of 10 μm below the processed surface was approximately 45 nm. 5. The deformed layer exhibited a mechanical gradient distribution where the outer region exhibited a much higher hardness than the inner region of the deformed layer and matrix, in which the hardness in the deformed layer ranged from 5.57 GPa near the surface to 4.2 GPa near the matrix. A strong strain hardening effect was observed in the deformed layer. The strengthening mechanism can be ascribed to grain boundary strengthening and dislocation strengthening.
4. Conclusions In this work, the ultrasonic surface rolling process was performed on a selective laser melted Ti6Al4V alloy, leading to a deformed layer with a gradient structure and significantly improved hardness. The main conclusions are summarized as follows:
Author contributions The author Zhen Wang did the experiment and wrote the article; the authors Chaofeng Gao and Zhongqiang Liu collected data; the author Zhi Wang improved the manuscript language; the authors Xiao Liu, Kaho Wong and Zhaoyao Zhou gave technical supports; The author Zhiyu Xiao provided helpful discussions and modification on the article.
1. A gradient structure deformed layer with a thickness of 300 μm was obtained when selective laser melted Ti6Al4V alloy was subjected to the USRP. From the matrix to the surface, the lamellar structure was broken, and both the average thickness and aspect ratio decreased. The gradient structure consisted of high-density dislocations, ultra fine lamellar grains, ultrafine equiaxed grains and equiaxed nanograins.
Declaration of competing interest The authors declare no competing financial interests. 7
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Acknowledgments
[22] G. Liu, J. Lu, K. Lu, Surface nanocrystallization of 316L stainless steel induced by ultrasonic shot peening, Mater. Sci. Eng. A 286 (2000) 91–95, https://doi.org/ 10.1016/S0921-5093(00)00686-9. [23] N.R. Tao, M.L. Sui, J. Lu, K. Lua, Surface nanocrystallization of iron induced by ultrasonic shot peening, Nanostructured Mater. 11 (1999) 433–440, https://doi. org/10.1016/S0965-9773(99)00324-4. [24] S. Lou, Y. Li, L. Zhou, X. Nie, G. He, W. He, Surface nanocrystallization of metallic alloys with different stacking fault energy induced by laser shock processing, Mater. Des. 104 (2016) 320–326, https://doi.org/10.1016/j.matdes.2016.05.028. [25] A.I. Dekhtyar, B.N. Mordyuk, D.G. Savvakin, V.I. Bondarchuk, I.V. Moiseeva, N. I. Khripta, Enhanced fatigue behavior of powder metallurgy Ti-6Al-4V alloy by applying ultrasonic impact treatment, Mater. Sci. Eng. A 641 (2015) 348–359, https://doi.org/10.1016/j.msea.2015.06.072. [26] B.N. Mordyuk, M.O. Iefimov, G.I. Prokopenko, T.V. Golub, M.I. Danylenko, Structure, microhardness and damping characteristics of Al matrix composite reinforced with AlCuFe or Ti using ultrasonic impact peening, Surf. Coat. Technol. 204 (2010) 1590–1598, https://doi.org/10.1016/j.surfcoat.2009.10.009. [27] B.N. Mordyuk, O.P. Karasevskaya, G.I. Prokopenko, Structurally induced enhancement in corrosion resistance of Zr-2.5%Nb alloy in saline solution by applying ultrasonic impact peening, Mater. Sci. Eng. A 559 (2013) 453–461, https://doi.org/10.1016/j.msea.2012.08.125. [28] Z. Wang, Z. Xiao, C. Huang, L. Wen, W. Zhang, Influence of ultrasonic surface rolling on microstructure and wear behavior of selective laser melted Ti-6Al-4V alloy, Materials 10 (2017) 1203, https://doi.org/10.3390/ma10101203. [29] G. Li, S.G. Qu, Y.X. Pan, X.Q. Li, Effects of the different frequencies and loads of ultrasonic surface rolling on surface mechanical properties and fretting wear resistance of HIP Ti–6Al–4V alloy, Appl. Surf. Sci. 389 (2016) 324–334, https:// doi.org/10.1016/j.apsusc.2016.07.120. [30] L. Chen, X. Ren, W. Zhou, Z. Tong, S. Adu-Gyamfi, Y. Ye, Y. Ren, Evolution of microstructure and grain refinement mechanism of pure nickel induced by laser shock peening, Mater. Sci. Eng. A 728 (2018) 20–29, https://doi.org/10.1016/j. msea.2018.04.105. [31] D. Liu, D. Liu, X. Zhang, C. Liu, N. Ao, Surface nanocrystallization of 17-4 precipitation-hardening stainless steel subjected to ultrasonic surface rolling process, Mater. Sci. Eng. A 726 (2018) 69–81, https://doi.org/10.1016/j. msea.2018.04.033. [32] N. Ao, D. Liu, X. Xu, X. Zhang, D. Liu, Gradient nanostructure evolution and phase transformation of α phase in Ti-6Al-4V alloy induced by ultrasonic surface rolling process, Mater. Sci. Eng. A 742 (2018) 820–834, https://doi.org/10.1016/j. msea.2018.10.098. [33] L.H. Su, C. Lu, L.Z. He, L.C. Zhang, P. Guagliardo, A.K. Tieu, S.N. Samarin, J. F. Williams, H.J. Li, Study of vacancy-type defects by positron annihilation in ultrafine-grained aluminum severely deformed at room and cryogenic temperatures, Acta Mater. 60 (2012) 4218–4228, https://doi.org/10.1016/j. actamat.2012.04.003. [34] M. Seifi, A. Salem, J. Beuth, O. Harrysson, J.J. Lewandowski, Overview of materials qualification needs for metal additive manufacturing, J. Occup. Med. (2016) 747–764, https://doi.org/10.1007/s11837-015-1810-0. [35] M. Simonelli, Y.Y. Tse, C. Tuck, On the texture formation of selective laser melted Ti-6Al-4V, Metall. Mater. Trans. A (2014), https://doi.org/10.1007/s11661-0142218-0. [36] C. Yang, J.A. Lin, Y.F. Ding, W.W. Zhang, Y.Y. Li, Z.Q. Fu, F. Chen, E.J. Lavernia, Texture evolution and mechanical behavior of commercially pure Ti processed via pulsed electric current treatment, J. Mater. Sci. 51 (2016) 10608–10619, https:// doi.org/10.1007/s10853-016-0282-0. [37] L. Wang, L. Xie, Y. Lv, L.C. Zhang, L. Chen, Q. Meng, J. Qu, D. Zhang, W. Lu, Microstructure evolution and superelastic behavior in Ti-35Nb-2Ta-3Zr alloy processed by friction stir processing, Acta Mater. 131 (2017) 499–510, https://doi. org/10.1016/j.actamat.2017.03.079. [38] Z. Liu, P. Li, L. Xiong, T. Liu, L. He, High-temperature tensile deformation behavior and microstructure evolution of Ti55 titanium alloy, Mater. Sci. Eng. A 680 (2017) 259–269, https://doi.org/10.1016/j.msea.2016.10.095. [39] Z. Ji, H. Yang, H. Li, Contributions of microstructural features to the integrated hardness of TA15 titanium alloy, Mater. Sci. Eng. A 628 (2015) 358–365, https:// doi.org/10.1016/j.msea.2014.12.087. [40] C.Y. Hyun, J.H. Lee, H.K. Kim, Microstructures and mechanical properties of ultrafine grained pure Ti produced by severe plastic deformation, Res. Chem. Intermed. 36 (2010) 629–638, https://doi.org/10.1007/s11164-010-0198-2. [41] P. Luo, D.T. McDonald, W. Xu, S. Palanisamy, M.S. Dargusch, K. Xia, A modified Hall-Petch relationship in ultrafine-grained titanium recycled from chips by equal channel angular pressing, Scr. Mater. 66 (2012) 785–788, https://doi.org/ 10.1016/j.scriptamat.2012.02.008. [42] B.J. Hayes, B.W. Martin, B. Welk, S.J. Kuhr, T.K. Ales, D.A. Brice, I. Ghamarian, A. H. Baker, C.V. Haden, D.G. Harlow, H.L. Fraser, P.C. Collins, Predicting tensile properties of Ti-6Al-4V produced via directed energy deposition, Acta Mater. 133 (2017) 120–133, https://doi.org/10.1016/j.actamat.2017.05.025. [43] R. Kapoor, N. Kumar, R.S. Mishra, C.S. Huskamp, K.K. Sankaran, Influence of fraction of high angle boundaries on the mechanical behavior of an ultrafine grained Al-Mg alloy, Mater. Sci. Eng. A 527 (2010) 5246–5254, https://doi.org/ 10.1016/j.msea.2010.04.086. [44] J. Moon, S. Kim, J. il Jang, J. Lee, C. Lee, Orowan strengthening effect on the nanoindentation hardness of the ferrite matrix in microalloyed steels, Mater. Sci. Eng. A 487 (2008) 552–557, https://doi.org/10.1016/j.msea.2007.10.046.
This work was financially supported by the National Natural Science Foundation of China (Grant No. 51627805), the Natural Science Foun dation Project of Guangdong Province (Grant No. 2015A030312003), and the Science and Technology Research Projects of Guangdong Province (Grant No. 2014B010129003, 2015B020238008, 2016B090931006, 2017B090901025). References [1] C.T. dos Santos, C. Barbosa, M.J. Monteiro, I.C. Abud, I.M.V. Caminha, C.R. M. Roesler, Characterization of the fretting corrosion behavior, surface and debris from head-taper interface of two different modular hip prostheses, J. Mech. Behav. Biomed. Mater. 62 (2016) 71–82, https://doi.org/10.1016/j.jmbbm.2016.04.036. [2] D. Gu, Y. Hagedorn, W. Meiners, G. Meng, Densification behavior , microstructure evolution , and wear performance of selective laser melting processed commercially pure titanium, Acta Mater. 60 (2012) 3849–3860, https://doi.org/ 10.1016/j.actamat.2012.04.006. [3] G. Kasperovich, J. Hausmann, Improvement of fatigue resistance and ductility of TiAl6V4 processed by selective laser melting, J. Mater. Process. Technol. 220 (2015) 202–214, https://doi.org/10.1016/j.jmatprotec.2015.01.025. [4] Y.H. Lu, Z.Y. Xiao, L. Hu, F. Luo, Y.B. Wu, D.H. Ni, Ultra-high cycle fatigue behaviour of warm compaction Fe-Cu-Ni-Mo-C sintered material, Mater. Des. 55 (2014) 758–763, https://doi.org/10.1016/j.matdes.2013.10.049. [5] I. Serrano-Munoz, J.Y. Buffiere, R. Mokso, C. Verdu, Y. Nadot, Location, location &size: defects close to surfaces dominate fatigue crack initiation, Sci. Rep. 7 (2017), 45239, https://doi.org/10.1038/srep45239. [6] R. Sonntag, J. Reinders, J. Gibmeier, J.P. Kretzer, Fatigue performance of medical Ti6Al4V alloy after mechanical surface treatments, PLoS One 10 (2015), e0121963, https://doi.org/10.1371/journal.pone.0121963. [7] L.C. Zhang, L.Y. Chen, A review on biomedical titanium alloys: recent progress and prospect, Adv. Eng. Mater. 21 (2019) 1–29, https://doi.org/10.1002/ adem.201801215. [8] B. Subramanian, C.V. Muraleedharan, R. Ananthakumar, M. Jayachandran, A comparative study of titanium nitride (TiN), titanium oxy nitride (TiON) and titanium aluminum nitride (TiAlN), as surface coatings for bio implants, Surf. Coat. Technol. 205 (2011) 5014–5020, https://doi.org/10.1016/j.surfcoat.2011.05.004. [9] S. Kumar, T.S.N.S. Narayanan, S. Ganesh Sundara Raman, S.K. Seshadri, Surface modification of CP-Ti to improve the fretting-corrosion resistance: thermal oxidation vs. anodizing, Mater. Sci. Eng. C 30 (2010) 921–927, https://doi.org/ 10.1016/j.msec.2010.03.024. [10] A. Bloyce, P.-Y. Qi, H. Dong, T. Bell, Surface modification of titanium alloys for combined improvements in corrosion and wear resistance, Surf. Coat. Technol. 107 (1998) 125–132, https://doi.org/10.1016/S0257-8972(98)00580-5. [11] M. Niinomi, Recent metallic materials for biomedical applications, Metall. Mater. Trans. A 33 (2002) 477–486, https://doi.org/10.1007/s11661-002-0109-2. [12] X.H. Chen, J. Lu, L. Lu, K. Lu, Tensile properties of a nanocrystalline 316L austenitic stainless steel, Scr. Mater. 52 (2005) 1039–1044, https://doi.org/ 10.1016/J.SCRIPTAMAT.2005.01.023. [13] H. Gleiter, Nanostructured materials: basic concepts and microstructure, Acta Mater. 48 (2000) 1–29, https://doi.org/10.1016/S1359-6454(99)00285-2. [14] R.K. Nalla, I. Altenberger, U. Noster, G.Y. Liu, B. Scholtes, R.O. Ritchie, On the influence of mechanical surface treatments-deep rolling and laser shock peening-on the fatigue behavior of Ti-6Al-4V at ambient and elevated temperatures, Mater. Sci. Eng. A 355 (2003) 216–230, https://doi.org/10.1016/S0921-5093(03)000698. [15] X. Liu, Z.y. Xiao, H.j. Guan, W. Zhang, F.l. Li, Friction and wear behaviours of surface densified powder metallurgy Fe–2Cu–0.6C material, Powder Metall. 59 (2016) 329–334, https://doi.org/10.1080/00325899.2016.1242880. [16] K.Y. Zhu, A. Vassel, F. Brisset, K. Lu, J. Lu, Nanostructure formation mechanism of α-titanium using SMAT, Acta Mater. 52 (2004) 4101–4110, https://doi.org/ 10.1016/j.actamat.2004.05.023. [17] L. Jin, W. Cui, X. Song, L. Zhou, The formation mechanisms of surface nanocrystallites in β-type biomedical TiNbZrFe alloy by surface mechanical attrition treatment, Appl. Surf. Sci. 347 (2015) 553–560, https://doi.org/10.1016/ j.apsusc.2015.04.137. [18] X.C. Liu, H.W. Zhang, K. Lu, Formation of nano-laminated structure in nickel by means of surface mechanical grinding treatment, Acta Mater. 96 (2015) 24–36, https://doi.org/10.1016/j.actamat.2015.06.014. [19] W.L. Li, N.R. Tao, K. Lu, Fabrication of a gradient nano-micro-structured surface layer on bulk copper by means of a surface mechanical grinding treatment, Scr. Mater. 59 (2008) 546–549, https://doi.org/10.1016/j.scriptamat.2008.05.003. [20] Y.G. Liu, M.Q. Li, H.J. Liu, Surface nanocrystallization and gradient structure developed in the bulk TC4 alloy processed by shot peening, J. Alloy. Comp. 685 (2016) 186–193, https://doi.org/10.1016/j.jallcom.2016.05.295. [21] H. Li, Y. Liu, M. Li, H. Liu, The gradient crystalline structure and microhardness in the treated layer of TC17 via high energy shot peening, Appl. Surf. Sci. 357 (2015) 197–203, https://doi.org/10.1016/j.apsusc.2015.09.013.
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