Microstructure evolution in Ti64 subjected to laser-assisted ultrasonic nanocrystal surface modification

Accepted Manuscript Microstructure Evolution in Ti64 Subjected to Laser-assisted Ultrasonic Nanocrystal Surface Modification

Jun Liu, Sergey Suslov, Zhencheng Ren, Yalin Dong, Chang Ye PII:

S0890-6955(18)30294-3

DOI:

10.1016/j.ijmachtools.2018.09.005

Reference:

MTM 3374

To appear in:

International Journal of Machine Tools and Manufacture

Received Date:

25 July 2018

Accepted Date:

23 September 2018

Please cite this article as: Jun Liu, Sergey Suslov, Zhencheng Ren, Yalin Dong, Chang Ye, Microstructure Evolution in Ti64 Subjected to Laser-assisted Ultrasonic Nanocrystal Surface Modification, International Journal of Machine Tools and Manufacture (2018), doi: 10.1016/j. ijmachtools.2018.09.005

This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT

Microstructure Evolution in Ti64 Subjected to Laser-assisted Ultrasonic Nanocrystal Surface Modification Jun Liu1, Sergey Suslov2, Zhencheng Ren1, Yalin Dong1*, Chang Ye1* 1Department 2Qatar

of Mechanical Engineering, University of Akron, Akron, OH 44325, United States

Environment and Energy Research Institute (QEERI), Qatar Foundation, Doha, Qatar Corresponding author e-mail address: [email protected], [email protected]

Declaration of Interests: None

ACCEPTED MANUSCRIPT

Microstructure Evolution in Ti64 Subjected to Laser-assisted Ultrasonic Nanocrystal Surface Modification Abstract Surface severe plastic deformation (SSPD) can significantly improve the mechanical properties of metallic components by inducing surface nanocrystallization and beneficial compressive residual stresses. The effectiveness of the SSPD processes is significantly dependent on the plasticity of the target metals. Here, we report an innovative surface thermomechanical process called laserassisted ultrasonic nanocrystal surface modification (LA-UNSM) that integrates localized laser heating with high strain rate plastic deformation. The laser beam locally heats the target metal and increases the local plasticity, making the SSPD treatment more effective. After LA-UNSM, a microstructure featuring a nanocrystalline layer embedded with nanoscale precipitates was achieved in Ti64, resulting in an unprecedented 75.2% increase in hardness. After LA-UNSM processing, a 25-μm severe plastic deformation layer was produced that was 2.5 times thicker than that of the room-temperature UNSM-processed material. The grains at the top surface were refined down to 37 nm, indicating a similar degree of nanocrystallization to that produced by UNSM at room temperature. Nanoscale precipitate particles with diameters in the range of 5~21 nm were non-uniformly distributed in the nanocrystalline surface layer. These precipitates were produced through laser-assisted dynamic precipitation. The extremely high surface strength obtained for the Ti64 was attributed to the composite microstructure featured by nanoscale grains embedded with nanoscale precipitates and the work-hardening. Keywords: Laser heating; laser-assisted ultrasonic nanocrystal surface modification; nanocrystalline; precipitates; Ti64; dynamic precipitation 1. Introduction The majority of metallic component failures originate at the surface of the component. This makes surface severe plastic deformation (SSPD), a process that improves the mechanical properties of metallic components by inducing plastic deformation in the near-surface region, extremely important. Notable SSPD processes include laser shock peening [1–3], shot peening [4,5], ultrasonic peening [6,7], surface mechanical attrition treatment [8,9], and most recently developed ultrasonic nanocrystal surface modification (UNSM) [10,11]. These SSPD processes, by imposing microstructure changes and beneficial compressive residual stresses, can significantly improve the resistance of metallic components to fatigue, wear and corrosion. The effectiveness of SSPD processes are dependent on the plasticity of the processed metal in addition to the processing conditions. Many high-strength alloys, especially those having a hexagonal close-packed (hcp) crystal structure (e.g., titanium alloys), have low plasticity because they have fewer active slip systems. The processing of these metals requires aggressive conditions (e.g., high peening intensity), which could lead to surface/subsurface cracking. For example, to process some high-strength metals—such as spring steel or titanium and its alloys—a high ultrasonic vibration amplitude is needed, and this often results in surface fracture, leading to poor surface finish and ultimately to poor fatigue performance and low corrosion resistance. By increasing the deformation temperature, warm/hot working can effectively treat hard and brittle metals without cracking by virtue of the higher dislocation mobility and thus higher plasticity achieved as the temperature increases. To increase the plasticity of the metal, a number of surface 1

ACCEPTED MANUSCRIPT

thermomechanical treatments have been developed. For example, shot peening at the optimal temperature can increase the magnitude of the compressive residual stress in spring steel because of the decrease in flow stress [12]. By applying ultrasonic surface rolling at elevated temperatures (80°C, 120°C and 160°C), a thicker plastic deformation layer with lower hardness relative to its room-temperature counterpart was produced in the Ti64 [13]. In addition to improving plasticity, high temperature could also lead to beneficial microstructure changes. For example, high-density dislocations, nanocrystallization, and even amorphization were observed in the surface layer of H62 brass after warm laser shock peening (WLSP), leading to improvements in both tensile strength and ductility [14]. By increasing the sample temperature, WLSP can induce highly dense nanoscale precipitate particles in both aluminum [15] and carbon steel [16]. These nanoscale precipitates were generated as a result of thermally-assisted dynamic precipitation in WLSP. The abundance of potential nucleation sites and the high nucleation efficiency results in the highly dense nanoscale precipitates [17–20]. These precipitates increase the surface strength and the stability of the compressive residual stresses generated by LSP by impeding the movement of dislocations [15,16]. Unfortunately, high temperature bulk heating leads to sample distortion/oxidation, poor tool life, and high lubrication costs. In addition, it would be extremely difficult and energy-inefficient to heat large samples in the required manufacturing processes. Since SSPD processes involve treating only the surface of the components, we only need to improve the plasticity of the metal surface through localized heating. This can be done using a laser beam, which has been widely used in industry for localized heating with high efficiency [21–24]. While traditional warm/hot deformation processes heat the entire workpiece, a laser beam can rapidly heat a localized region to a very high temperature, which prevents the heating of the entire sample [25–29]. For example, laser-assisted machining has been used to soften the material using a laser beam to enable the machining of difficult-to-machine materials [30]. It is thus of great interest to integrate laser heating with the SSPD processes. Unfortunately, this approach has rarely been reported in the literature except for a study by Tian and Shin [31], in which the authors developed a laser-assisted burnishing (LAB) process that integrates laser heating with burnishing. The localized heating induced by the laser beam was reported to improve metal plasticity and make the burnishing treatment more effective. As compared to room-temperature burnishing, LAB results in higher hardness in the surface layer of the sample due to enhanced plastic deformation and work hardening. In this study, we integrate laser heating with the recently developed SSPD process UNSM. It can significantly improve the wear and fatigue resistance of metallic materials by introducing plastic strain in the near surface region. In a UNSM process, a tungsten carbide tip strikes the surface of the workpiece to induce severe plastic deformation (SPD) [32]. By superimposing a variable static load (typically 10 to 50 N) and dynamic amplitude (typically 8 to 50 μm) on the tip, the intensity of the strike can be precisely controlled as compared with a hand-held ultrasonic impact technique (UIT) [33]. The UNSM scan path and the distance between neighbor scans can also be precisely controlled using computer numeric control (CNC) codes, significantly improving the reproducibility of UNSM as compared with shot peening (SP) and surface mechanical attrition treatment (SMAT) [9,34,35]. Through UNSM treatment, compressive residual stress, nanograins and high-density dislocation can be introduced into the workpiece surface and subsurface, thereby enhancing the surface strength, tribological performance, and fatigue properties of the metallic material. This technique has been widely 2

ACCEPTED MANUSCRIPT

applied to process stainless steels [36], aluminum alloys [37], magnesium alloys [32], Cu-based alloys [38], and titanium and its alloys [39,40]. Ti64 alloy is chosen as the material in this study, since Ti64 has become the workhorse titanium alloy for biomedical, marine, automotive, and aerospace applications thanks to its high strength, low weight ratio, excellent fatigue performance and superior corrosion resistance. The poor antifriction characteristics and low hardness of Ti64, however, have hindered more extensive application of this alloy. In addition, Ti64 has a high work hardening tendency, which makes it difficult to further improve the surface strength using traditional surface treatments [41]. Thus, it is highly desirable to enlarge the thickness of the strengthening layer in Ti64 while further improving the surface strength of the alloy. Previous studies have demonstrated that the mechanical properties of the alloy can be tailored using heat treatments or thermomechanical processing by appropriately modifying the alloy’s microstructure [42,43], particularly in terms of alpha (α) and beta (β) phases. For example, cooling at a high rate from the α+β region or above the β-transus temperature causes the β-Ti to transform into a martensitic phase with two possible forms: α’ (hcp) and α’’ (orthorhombic) [43,44]. By over-aging, a Ti64 alloy is prone to precipitation of α2 (Ti3Al) [45,46]. Yu et al. reported secondary α precipitation after isothermal compression at temperatures of 910~980°C [47]. Therefore, LA-UNSM has potential as a thermomechanical process to customize the microstructure of the alloy and thereby improve its mechanical properties. Additionally, in LA-UNSM, SPD is introduced into the material by mechanical impact of the tip on the target surface. The tip is made of tungsten carbide, which has a melting point over 2785 °C. Meanwhile, thermal energy is directly delivered onto the same region by the laser beam. Therefore, the processing temperature of LA-UNSM is highly adjustable with wider temperature range as compared with WLSP. The laser beam of LA-UNSM can bypass the challenge of heating the bulk specimen and thereby be applied to process products with complex shapes. In the present study, we investigate the microstructure evolution and its effect on mechanical properties in Ti64 subjected to LA-UNSM. The microstructure for the Ti64 samples after UNSM and LA-UNSM processing was characterized using X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM), and scanning TEM (STEM). The nanoscale precipitates were characterized using high-resolution transmission electron microscopy (HR-TEM). Thermographic images of the workpiece during LA-UNSM processing were recorded. The microhardness test was carried out to demonstrate the substantial improvement in surface strength resulting from LA-UNSM processing. It was observed that a nanocrystalline microstructure embedded with nanoscale precipitates was achieved in Ti64 after LA-UNSM treatment. This unique microstructure results in extremely high hardness. 2. Experimental work 2.1. LA-UNSM process Test specimens (14 mm × 14 mm × 3 mm) were cut from a Ti64 sheet having a chemical composition of 0.08 wt% carbon, 5.50–6.75 wt% aluminum, 3.5–4.5 wt% vanadium, 0.4 wt% max. iron, 88.10–90.92 wt% titanium and 0–0.3% other components. In a traditional UNSM process, a tungsten carbide ball (with a diameter of 2.38 mm) attached to an ultrasonic device scans over the material surface while striking it at high frequency (20 kHz). The overlapping of the mechanical impacts generates plastic strain at the material surface and leads to surface plastic deformation. The processing parameters in the UNSM treatment include the static load, the 3

ACCEPTED MANUSCRIPT

distance between neighbor scans (interval), and the scanning speed, all of which can be precisely controlled using an integrated CNC machining system. In this work, the UNSM processing of the Ti64 alloy was performed using the parameters listed in Table 1. These parameters were selected based on previous studies [39,40]. Table 1. UNSM processing parameters Amplitude Condition

(μm)

Static load (N)

UNSM

24

50

Scanning

Interval

(mm/min)

(μm)

Tip diameter (mm)

1000

10

2.38

An overview of LA-UNSM process is shown in Fig. 1a. During LA-UNSM, the workpiece surface is locally heated using a continuous laser beam and then immediately processed by a UNSM tool. The working temperatures for LA-UNSM are controlled using a neodymium-doped yttrium aluminum garnet (Nd: YAG) continuous wave fiber laser (a 1064-nm, 200-W, model YLR-200MM-AC-Y11 laser from IPG Photonics). The optical head of the laser is attached to an ultrasonic unit (Fig. 1b) so that the two components will move together; this synchronized movement ensures that simultaneous laser heating and ultrasonic striking will occur at the same location. The unit delivers a laser beam with a diameter of 0.74 mm onto the target sample through a feeding fiber (an LC-8 with a 50-μm core diameter), supplying the necessary flexibility to adjust the relative position of the laser beam with respect to the sample and the UNSM tip. The test setup for the LAUNSM process is shown in the schematic diagram in Fig. 1c.

Figure 1. LA-UNSM test setup: (a) Overview of the LA-UNSM system, (b) magnification of the sample during LA-UNSM processing from the view b in Fig. 1a, and (c) schematic diagram of the LA-UNSM process. 4

ACCEPTED MANUSCRIPT

2.2. Material characterization Material characterization for the Ti64 samples was conducted using the following equipment and procedures: 

SEM characterization: For cross-sectional studies, specimens of the as-received (AR), UNSM-processed, and LA-UNSM-processed Ti64 samples were obtained from the crosssection, polished by standard metallographic techniques to a mirror-like finish, followed by etching using Kroll’s reagent (a mixture of distilled water, nitric acid, and hydrofluoric acid) for 15 s. The surface morphology of the top surfaces and cross sections was investigated by SEM using a Tescan Lyra 3 system operated at 30 keV, and the chemical composition was confirmed using energy-dispersive X-ray spectroscopy (EDS).



XRD characterization: XRD patterns were collected using a Rigaku Ultima IV diffractometer with a Cu Kα-radiation source operated at 45 kV and 35mA at a scan speed of 1.0°/min.



TEM and STEM characterization. The TEM samples were prepared using the focused ion beam (FIB) lift-out method in an FEI Versa 3D LoVac FIB-SEM Dual Beam system equipped with an FEI EasyLift system. The TEM and STEM testing was carried out using an FEI Talos F200X TEM working at 200 kV.



Microhardness: The microhardness change for all samples was measured using a Wilson Tukon 2100 system with a Vickers indenter under a 50 g load and a 10 s dwell time. The average of five measurements was reported.



Temperature: Thermal imaging was performed using an infrared (IR) camera (FLIR T650sc, FLIR®). The images were analyzed using FLIR ResearchIR Max 4 software (from FLIR Systems, Inc.). To calculate the temperature of the workpiece, it is necessary to obtain the emissivity of the Ti64 alloy. In this study, the temperature was measured using a thermocouple (4004 Traceable® Big-Digit Type K), and IR images were collected at the same time. The arithmetic mean of 0.529 for the calibration results was taken as the emissivity for IR camera imaging.

3. Results 3.1. Optimal laser power for LA-UNSM analyzed by surface hardness Material strength is largely affected by microstructure such as grain size, dislocation density and distribution as well as precipitate size and density. Since processing temperature significantly affects dislocation movement, grain growth and precipitation, the optimal laser power for LAUNSM processing is material dependent. To find the optimal laser power to use for processing Ti64, LA-UNSM was performed with a laser power from 12 W to 68 W. The hardness is plotted as a function of laser power as shown in Fig. 2a. It can be observed that UNSM processing without a laser significantly increased (by 43.5%) the hardness of the AR Ti64 due to cold working. After LA-UNSM processing, the material hardness first decreases then gradually increases as laser power increases from 12 W to 56 W. The hardness of the sample after LA-UNSM at 12 W is lower than that after room temperature UNSM. Above 12 W, the hardness of the LA-UNSM processed samples becomes higher than that of the UNSM-processed sample. When the laser power reaches 5

ACCEPTED MANUSCRIPT

68 W, a drop in hardness is observed. Therefore, it can be concluded that 56 W is the optimal laser power for LA-UNSM processing of Ti64 alloy from the perspective of surface strengthening, which gives rise to an increase of approximately 22.2% in hardness as compared with that of room temperature UNSM (and an increase in hardness of 75.2% as compared with that of the AR Ti64 sample), indicating a remarkable improvement beyond the limit of cold working induced by UNSM. The following studies were carried out using this optimal laser power.

Figure 2. (a) Hardness as a function of laser power. AR refers to the Ti64 alloy in the as-received status (without UNSM or laser heating treatment). (b) Surface microhardness for Ti64 alloy: AR, laser scanning using 56 W power (AR-56W), UNSM alone, UNSM followed by laser scanning using 56 W power (UNSM-56W), and LA-UNSM (56 W) processing. 3.2. Verification of the coupling effects of laser heating and UNSM processing The extra hardness that beyond cold working of the LA-UNSM Ti64 results from the coupling of the high strain rate deformation and laser heating. To verify this, two additional experiments were performed: laser scanning of AR Ti64 at 56 W power (AR-56W) and UNSM treatment followed by laser scanning at 56 W (UNSM-56W). The resultant hardness is shown in Fig. 2b. It can be seen that the AR-56W shows a modest increase in hardness as compared with that of the AR sample; however, the hardness of the AR-56W is much lower than that of the LA-UNSM, indicating that the increase in hardness was not caused by laser heating alone. The UNSM-56W sample shows a slight decrease in hardness as compared to the UNSM sample (cold working alone). Therefore, it is confirmed that the hardness improvement of the LA-UNSM Ti64 results from concurrent thermal and mechanical effects that can only be achieved by simultaneous laser irradiation and UNSM processing. 3.3. LA-UNSM working temperature The temperature distribution caused by laser heating is important because temperature plays a key role in the microstructure evolution of a material. Fig. 3a is a digital image demonstrating the corresponding configuration of the tip and workpiece, and Figs. 3b to 3f show IR camera recordings of the temperature distribution across the surface of the material. It can be seen that the tip was heated first (where the maximum temperature was approximately 190°C), while the 6

ACCEPTED MANUSCRIPT

average sample temperature was lower than 100°C (Fig. 3b). After approximately 12 minutes, the measured temperature of the configuration, including the tip and the sample, showed a significant increase. To better observe the temperature distribution on the sample surface, temperatures below 100°C and above 300°C are filtered out as shown in Fig. 3d. The presence of temperatures higher than 300°C, indicated by the green arrow in Figs. 3d and 3f, suggests an intensive and local heating area that was confined to the contact area between the small tip and the workpiece. A thermographic image collected at the end of processing is shown in Fig. 3e, where temperatures of the entire configuration were in the range of 24.4°C~418.8°C. After filtering out the temperatures below 100°C and above 250°C (Fig. 3f), it can be observed that the high-temperature field was still confined to the small tip/workpiece contact region, while the rest of the sample was in the range of 100°C~250°C, implying a local heating caused by laser radiation throughout the whole LA-UNSM processing.

Figure 3. LA-UNSM processing: (a) Digital images of the configuration for LA-UNSM processing. IR camera image showing the workpiece temperature at (b) the beginning of the LAUNSM processing when the laser power was 56 W, (c) after 12 minutes of processing, 7

ACCEPTED MANUSCRIPT

(d) after filtering out temperatures below 100°C and above 300°C, (e) at the end of processing, and (f) after filtering out temperatures below 100°C and above 250°C. 3.4. Phase analysis by XRD The XRD curves for the UNSM and LA-UNSM Ti64 alloys are plotted in Fig. 4a. For detailed phase analysis, enlarged sections of the XRD spectra of all the presented reflections are shown in Figs. 4b–4f. It can be seen that nearly all the reflections of the LA-UNSM Ti64 present a slight shift to a higher angle as compared to that of the UNSM Ti64, indicating an increase in the magnitude of the compressive residual stress (Figs. 4b, 4c, 4e, and 4f). For cases where a higher laser power was used (above 23 W), the full width at half maximum (FWHM) values of the peaks are identical to that of the UNSM at some reflections planes, such as α(002), α(102), α(103) and α(201). Of particular interest is the presence of a new phase when the laser power is higher than 12 W, as is evident in Fig. 4d. The new phase has a peak region at around 43.5° (2θ), which very possibly is α2 (Ti3Al) phase. Furthermore, from 23 W to 56 W, the intensity of the new phase increases with the laser power, showing a clear correlation with the working temperature. However, a further increase in the laser power to 68 W leads to weakening of the intensity of the new phase. It is noted that the trend for the intensity of the new phase is nearly identical to the trend for the microhardness, as is evident in Fig. 2a. Thus, it is possible that the presence of the new precipitate phase plays a key role in the strengthening of the LA-UNSM-processed Ti64 alloys.

Figure 4. XRD spectra of UNSM and LA-UNSM Ti64 alloy workpieces: (a) XRD spectra under different laser powers. XRD spectra of (b) α(100) and (c) α(002) and α(101). (d) XRD spectra at around 43.5°. XRD spectra of (e) α(102) and α(110) and (f) α(103) and α(201). 8

ACCEPTED MANUSCRIPT

Fig. 5 shows the XRD patterns of the Ti64 alloy in the AR, UNSM and LA-UNSM (56 W) samples. In the AR Ti64, the diffraction pattern indicates a combination of hexagonal close packed α-Ti and body-centered cubic β-Ti. Titanium in β phase can be clearly characterized by the β (110) reflection. After the UNSM treatment, remarkable changes were introduced to the microstructure. Significant grain refinement, as well as heavy lattice distortions, are evidenced by a noticeable broadening in FWHM of the XRD peaks. The β-Ti reflection lost its intensity dramatically, revealing a β-to-α transformation induced by UNSM treatment, which was also reported by Vasylyev [7] for Ti64 alloy after UIT. After LA-UNSM treatment, all the α-Ti reflections experienced a significant loss of intensity. The β-Ti phase cannot be observed, while the new phase was observed at around 43.5° (2θ). The FWHM values of all the α-Ti reflections are listed in the Table 2. It is seen that the FWHM values of α(100), α(002), and α(103) of the UNSM treated sample are larger than that of the LA-UNSM treated sample. However, the reflections of α(101), α(102), α(110), and α(112) of the UNSM sample demonstrate lower FWHM values as compared to that of the LA-UNSM sample. Thus, it is concluded that the profiles of all the reflections in the LA-UNSM sample are comparable to those of the UNSM sample, implying similar degrees of refinement and work hardening effect.

Figure 5. XRD patterns of the AR, UNSM, and LA-UNSM Ti64 alloy workpieces: (a) XRD spectra at 56 W; (b) high magnification image of XRD spectra of α(100), α(002), β(110) and α(101); (c) XRD spectra at around 43.5°; high magnification images of XRD spectra of (d) α(102) and α(110) and (e) α(103), α(200), α(112), and α(201).

Table 2. Comparison of full width at half maximum (FWHM) of all α-Ti reflections of the AR, UNSM and LA-UNSM Ti64 alloy at 56 W

9

ACCEPTED MANUSCRIPT

Lattice plane

AR

UNSM

LA-UNSM (56W)

(100)

0.220

0.615

0.587

(002)

0.190

0.560

0.452

(101)

0.236

0.566

0.601

(102)

0.246

0.647

0.663

(110)

0.292

0.790

0.864

(103)

0.254

0.768

0.762

(200)

0.295

--

--

(112)

0.294

0.767

0.820

(201)

0.322

--

--

3.5. Microstructure observed by SEM  The LA-UNSM processing involves elevated temperature and thereby possibly produces metal oxides and/or nitrides when processes metallic materials. However, the resultant oxide and/or nitride layer should be thin and play a trivial role on surface hardness. This is because a LA-UNSM treatment was less than 25 minutes, which is not sufficient long enough to allow oxygen and nitrogen penetrate into the material to form a thick oxide or nitride layer. To determine the surface strength attributed to microstructure evolution, all samples surfaces were polished using 3 μm diamond suspension to remove the possible oxide or nitride layer before the hardness tests. We performed surface SEM characterizations and elemental mappings on the UNSM-treated and LAUNSM-treated (56W) Ti64 surfaces. As shown in Fig. 6, several parallel bands produced by the UNSM treatment can be observed. This is because during a UNSM process, the tip keeps in contact with the sample surface due to the high frequency vibration, which is analogous to a “plowing” process, leading to the presence of the regular grooves. The SEM-EDS elemental mapping images in Fig. 6b~e reveal the presences of O, Al, Ti, and V after UNSM treatment. It is found that the elements are uniformly distributed corresponding to the SEM image (Fig. 6a). However, the weight percentage (wt. %) of O is only 1.2 (Fig. 6f) and such small amount of O cannot affect the materials surface strength significantly. In LA-UNSM, the regular bands cannot be observed as shown in Fig. 7a due to the thermal-induced plasticity enhancement of the material. Fig. 7b~e demonstrate the SEM-EDS elemental mapping images of O, Al, Ti, and V after LA-UNSM treatment. The weight percentage of O is only 3.52, indicating the successful removal of the oxide 10

ACCEPTED MANUSCRIPT

layer. Therefore, it is safely to rule out the influence of oxide or nitride on surface hardness of the UNSM-treated and LA-UNSM-treated Ti64.

Figure 6. Surface SEM image of the UNSM-treated Ti64 alloy (a) and its elemental mapping of (b) O; (c) Al; (d) Ti; (e) V. (f) Quantification of the elements in wt. %.

11

ACCEPTED MANUSCRIPT

Figure 7. Surface SEM image of the LA-UNSM-treated Ti64 alloy (a) and its elemental mapping of (b) O; (c) Al; (d) Ti; (e) V. (f) Quantification of the elements in wt. %. Fig. 8 shows the cross-sectional SEM images of the AR, UNSM processed, and LA-UNSM processed Ti64 samples. The AR Ti64 alloy (Fig. 8a) normally consists of equiaxed α grains lying within the β matrix, demonstrating a uniform microstructure before UNSM treatment. The microstructure after UNSM is shown in Fig. 8b. It can be seen that the original β grains evolved into compressed and elongated grains in the direction parallel to the specimen surface, presenting an approximately 10-μm SPD zone caused by ultrasonic striking during UNSM. LA-UNSM also introduced a layer of SPD as shown in Fig. 8c. The morphology of the compressed grains is magnified in the inset. Additionally, the SPD produced by LA-UNSM is substantially deeper than that of UNSM, which gives rise to a plastically deformed layer that is approximately 25 μm in depth.

12

ACCEPTED MANUSCRIPT

Figure 8. Cross-sectional SEM images, showing the microstructure of the Ti64 alloys: (a) AR, (b) UNSM-treated, and (c) LA-UNSM-treated alloy. 3.6. Microstructure observed by TEM, STEM and HRTEM The microstructure of the Ti64 alloy after UNSM at room temperature was further investigated using TEM to capture microstructural features that were not resolvable by SEM. Fig. 9 shows bright field (BF) images at the top surface of the workpiece. The dislocation structures cannot be clearly observed since the material was severely deformed [48]. The grains of the Ti64 alloy after UNSM treatment have an irregular shape and are remarkably refined to the nanometer range. The grain sizes are in the range of 25~94 nm (measured using the open source image processing program ImageJ). As observed in Fig. 9b, the microstructure is featured by wavy and poorlydefined boundaries. The poorly defined grain boundaries are an indication of significant internal stress, as suggested in other studies [49–51]. Moreover, the presence of high-density lattice defects results in a complex, non-uniform contrast microstructure in the Ti64 alloy, indicating a strong work hardening effect [43]. Hence, the microstructure confirms that UNSM can induce strong work hardening as well as nanocrystallization in the Ti64 alloy.

Figure 9. TEM micrograph of Ti64 alloy: (a) microstructure after UNSM treatment at room temperature and (b) magnification of area b. 13

ACCEPTED MANUSCRIPT

Figs. 10a to 10d show representative TEM micrographs of the Ti64 alloy after LA-UNSM processing. It can be seen that LA-UNSM allows for a distinctly globular structure of equiaxed grains with rather low defect density and well-distinguished boundaries at the very top layer, as shown in Fig. 10b. The grain size was measured, and it was found to be in the range of 37~77 nm, indicating no significant grain coarsening in comparison to that of the refined grains in the UNSM sample. As shown in Figs. 10c and 10d, extensive plastic strain and grain refinement below the equiaxed grain layer are evidenced by elongated nanograins with larger sizes as compared to the grains in the topmost layer (Fig. 10b). The morphology and grain size indicate the attenuated grain refinement resulting from the gradient nature of the plastic strain generated by LA-UNSM. We measured the grain size (indicated by the arrow in Fig. 10c): the long axis of the grain was 472 nm and the short axis was 87 nm. This microstructure proves that LA-UNSM can realize nanocrystallization of Ti64 alloy to the same degree as UNSM.

Figure. 10. A cross-sectional TEM micrograph of LA-UNSM Ti64 alloy: (a) subsurface microstructure of alloy treated by LA-UNSM. (b) A magnified TEM image shows equiaxed grains below the topmost surface in area b from Fig. 10a. TEM images showing the elongated grains in area c from Fig. 10a at a depth of approximately 1.8 μm below the surface: (c) a bright field (BF) image and (d) a dark field (DF) image. 14

ACCEPTED MANUSCRIPT

The STEM micrographs (Fig. 11) show the microstructure of the LA-UNSM-processed Ti64 at different levels of magnification. In addition to the significant grain refinement, high-density granular particles with much smaller sizes as compared to the nanograins can be observed in certain areas (Fig. 11c and 11d), identifying the presence of precipitates that cannot be observed in the UNSM-processed samples. This observation is consistent with the XRD results shown in Fig. 4 and Fig. 5, which suggest that a new second phase was produced after LA-UNSM treatment. In contrast, no evidence of room-temperature precipitation was observed in the UNSM-treated alloy. It is of particular interest to notice that the precipitates are not uniformly distributed in the matrix. Instead, the microstructure can be distinguished as nanograin zones (NZ) and precipitaterich zones (PZ). For example, in Fig. 11c and 11d, the NZs and PZs, as well as the boundaries between them (which are indicated by dashed lines), can be clearly observed. The nanograins can be observed between two PZs in the dark field TEM image (Fig. 11c). The PZ are dotted with high density granular particles, displaying a comparatively uniform gray color as shown in Fig. 11d. Therefore, as shown in Fig. 11a, the areas indicated by the yellow arrows are PZs, and the green arrows indicate the NZs. To clearly demonstrate the microstructure, a schematic of Fig. 11a is depicted as shown in Fig. 11b. Microstructural analysis of the LA-UNSM-processed Ti64 alloy revealed that the PZs interweave with the NZs to form a composite construction.

15

ACCEPTED MANUSCRIPT

Figure 11. Near surface region of the LA-UNSM-processed Ti64 alloy: (a) Cross-sectional BF STEM micrograph of the near surface and (b) schematic of the area shown in the micrograph. Higher magnification images of area c from Fig. 11a: (c) BF image and (d) DF image. In Fig. 12a, the distribution of the precipitates are characterized within a PZ. In this zone, a very dense network of precipitates that are homogeneously distributed was observed, as shown in Fig. 10b. The size of the precipitate, measured using ImageJ, was found to be in the range of 5 nm to 21 nm, with an average of 11 nm. Fig. 12a also reveals a large volume fraction of precipitates with a number density of > 36 per square micrometer. The density seems higher than that reported for conventional precipitation-hardened Ti64 alloy such as processed by over-aging [52]. To gain more insight on the precipitate structure, HRTEM image were obtained for a typical precipitate in the LA-UNSM sample. A high-resolution TEM image of the precipitate in the top surface indicated by an arrow in Fig. 12b is shown in Fig. 12c. It can be observed that the precipitate has a globular shape with a diameter of 11.5 nm. The corresponding diffraction pattern is given in Fig 12d. This pattern confirms the single crystalline nature of the precipitate with high symmetry cubic L12 structure. The precipitates reported by previous studies is α2 (Ti3Al) with hexagonal DO19 structure in Ti-Al alloys [53], which is conflict with our findings. A transformation of Ti3Al from DO19 (or a hexagonal close packed Ti) to L12 (or face centered cubic Ti) has been proved involving high pressure or high-strain rate plastic deformation [54–56]. The traditional precipitate of α2 with DO19 structure is normally produced through over ageing without sufficient plastic deformation [52]. Herein, we presume the precipitate produced in the LA-UNSM-treated samples is Ti3Al but with L12 structure. The cubic L12 Ti3Al results from the high-strain rate plastic deformation induced by UNSM combined with heat energy introduced by laser scanning. The transformation mechanism still need further investigation.

Figure. 12. PZ zone of the LA-UNSM-processed Ti64 alloy: (a) Cross-sectional STEM BF image, (b) HRTEM image, (c) magnified HRTEM image of the precipitate indicated by the arrow, and (d) electron diffraction pattern for the indicated precipitate.

16

ACCEPTED MANUSCRIPT

4. Discussion In this section, the mechanisms responsible for the microstructure features generated in Ti64 using LA-UNSM are discussed. 4.1. Nanograin/nanoprecipitate structure generated by LA-UNSM A unique nanograin/nanoprecipitate structure was produced by LA-UNSM in the near-surface region of Ti64 as shown in Fig. 11. The original grains were refined to nanoscale due to the plastic strain induced by UNSM [39,40]. Deformation-induced grain refinement and nanocrystallization of α- and β-Ti have been well discussed in previous studies [35,48,57,58]. Herein, we focus on the formation of the nanoscale precipitates. The UNSM and LA-UNSM Ti64 have similar grain size (UNSM: 25~94 nm; LA-UNSM: 37~77 nm), indicating a similar degree of refinement. The difference is that after the LA-UNSM treatment, nanoprecipitates were generated and were nonuniformly distributed in the nanograins. It seems that the nanoprecipitates were generated in the α-phase since Al is originally dispersed in the phase as the stabilizer. Accordingly, we propose the following mechanism for the generation of such unique nanostructure: assisted by the thermal energy from laser heating, the precipitates were nucleated at the defects (dislocations, twins, and triple junctions) induced by the plastic deformation and grew nanoscale precipitates. The process of nanocrystallization and precipitation of Ti64 during LA-UNSM is schematically illustrated in Fig. 13.

17

ACCEPTED MANUSCRIPT

Figure 13. Schematic of the microstructure evolution induced by LA-UNSM. For plastic deformation of Ti64 alloy, due to fewer slip systems in hcp Ti, twins are generated (Fig. 13b) and dominate the plastic deformation process [9,58]. Herein, the twin regions close to the top surface were fully developed into nanograins and nano-precipitates after LA-UNSM, which makes the twins are difficult to be captured using TEM images (Fig. 11). However, twinning deformation still plays a role in the plastic deformation of α-Ti at elevated working temperature [59–61]. At the same time, dislocations are triggered and accumulated in β-Ti. The Ti3Al precipitate in α-Ti mainly nucleate in regions close to twins, grain boundaries, and triple junctions as shown in Fig. 13c. With the superposition of subsequent strikes, more plastic deformation is introduced into the material. The accumulation and multiplication of the dislocations subdivide the grains into smaller pieces [9,35]. During this process, the subgrain boundaries and newly formed grain boundaries provide additional nucleation sites for the precipitates, which significantly increases its density as shown in Fig. 13d. These regions finally develop into PZs while other regions without nuclei are refined to nanograins as shown in Fig. 13e. 4.2. Nano-precipitates produced by dynamic precipitation The nanoscale precipitates were produced through laser-assisted dynamic precipitation (DP). The processing duration of the UNSM or the LA-UNSM was less than 25 min., and the maximum working temperature was below 500°C. In studies of the literature, precipitates in Ti64 were generated only after processing with temperatures above 500°C for several or even hundreds of hours [43,45–47]. Thus, the precipitate generated in the present study cannot be attributed to static over ageing. With the assistance of plastic deformation, the precipitation kinetic can be substantially enhanced (strain-induced precipitation). The defects generated by plastic deformation, including voids, dislocations, twins and grain boundaries, act as preferential nucleation sites for precipitation [62,63]. That is to say, nucleation primarily takes place in regions close to the defects [63]. Therefore, the density of available nucleation sites is proportional to the defect density. This explains why the high-density precipitate can be observed in short processing time. Additionally, the nucleation rate can be enhanced by the processing temperature during LA-UNSM, which can be mathematically described by an equation from Russell [64]: 𝑑𝑁 Δ𝐺 𝜏 ' ' = exp ( ‒ )exp ( ‒ )𝑁0, 𝑍 𝛽 𝑑𝑡 𝑘𝑇 𝑡

(1)

where N is the number of precipitates per unit volume and t is the time; thus, '

'

𝑑𝑁 𝑑𝑡

is the nucleation

rate. On the right-hand side of the equation, 𝑍 is Zeldovich’s factor, 𝛽 is the atomic impingement rate, 𝑇 is the processing temperature, and τ is the incubation period. The activation energy, Δ𝐺, has four components: Δ𝐺 =‒ Δ𝐺𝑐ℎ𝑒𝑚 + Δ𝐺𝜀 + 𝛾𝑖 ‒ Δ𝐺𝑑,

(2)

where Δ𝐺𝑐ℎ𝑒𝑚 is the chemical driving force, Δ𝐺𝜀 is the volume strain energy, 𝛾𝑖 is the interfacial free energy, and Δ𝐺𝑑 is the dislocation core energy. 18

ACCEPTED MANUSCRIPT

From Equation (1), we can see that the nucleation rate increases with the working temperature. During the LA-UNSM treatment, the laser beam heated the material to over 400°C (as shown in Fig. 3), which accelerates the nucleation rate. In addition, the nucleation rate is also determined by the activation energy, which is strongly affected by the chemical driving force as expressed in Equation (2) [17]. For LA-UNSM, the chemical driving force for nucleation was reduced by the elevated working temperature. As illustrated in Fig. 14a, the chemical driving force is the vertical difference (indicated by the red arrows) between the free energy of the matrix and nucleus. The nucleus is normally in a relatively stable state, which has a lower free energy, until reaching the equilibrium temperature 𝑇𝑒. Thus, the excess free energy in the matrix provides the chemical driving force for nucleation, and this energy is reduced by the elevated working temperature. This results in lower activation energy at a higher working temperature, as indicated by Equation (2), which then accelerates the nucleation rate, as calculated in Equation (1).

Figure 14. (a) The chemical driving force vs. the working temperature for LA-UNSM process. (b) Surface hardness for UNSM at different laser powers and the corresponding XRD peak intensity of the precipitate. 4.3. Enhanced hardening by nanoscale precipitates The presence of the precipitates lead to strengthening of the Ti64. The results of the hardness tests (shown in Figs. 2) clearly demonstrate extra hardening caused by the coupling of the mechanical and thermal treatments. However, annealing-induced recovery and recrystallization of materials normally lead to a decrease in strength. In previous research on processing Ti64 alloy at low temperatures (80°C, 120°C and 160°C), the thickness of the hardening layer increased due to a decrease of flow stress, while the absolute hardness was decreased in comparison to that of the room temperature processed alloy [11]. Although the working temperature in the LA-UNSM (56 W) was higher than that reported by Li et al. [11], the absolute value of the hardness is much higher than its room temperature counterpart (UNSM), which can be attributed to the presence of the precipitates. In LA-UNSM-processed Ti64 alloys, precipitation hardening essentially dominates the strengthening of the workpiece. As shown in Fig. 2 and Fig. 4d, the absence of precipitation in LA-UNSM at 12 W results in a decrease in hardness in comparison to that of UNSM. Starting at 19

ACCEPTED MANUSCRIPT

a power of 23 W, the presence of precipitates in the LA-UNSM workpiece compensates for the loss in hardness and gives rise to an additional increase in hardness; the hardness reaches the maximum value at 56 W. Meanwhile, the volume fraction of the precipitate increases with laser power. To better observe the correlation between the hardness and precipitation, the surface hardness for UNSM at different laser powers and the XRD peak intensity of the precipitate were replotted as shown in Fig. 14b. It can be seen that the hardness and peak intensity of the precipitate increase nearly linearly with the laser power from 12 W to 56 W, followed by an obvious drop for both at 68 W. Thus, the hardness increase is strongly associated with the density of the precipitates. The precipitates in the LA-UNSM-treated alloy could impede dislocation movement via the pinning effect. It is well established that the precipitation of a secondary phase within the lattice of a material will create physical blockades. These blockades act to halt dislocation movement, requiring greater energy or stress to overcome the barrier. Before dislocations cut through the precipitates, the dislocations must bow around these blockades, and this induces higher strength. 𝐺𝑏

This mechanism is called Orowan strengthening and can be mathematically described by 𝜏 = 𝐿 ‒ 2𝑟, where τ is the material strength, G is the shear modulus, b is the magnitude of the Burgers vector, L is the distance between pinning points, and r is the precipitate radius. This governing equation shows that for dislocation bowing, the strength is inversely proportional to the distance between the pinning points, L. In the present study, the density of the precipitate is very high as shown in Fig. 12a. That is to say, the average distance between pinning points is short. Thus, the strength of the LA-UNSM Ti64 alloy is high. Plus the nanograins produced by SPD, it is therefore concluded that the increase in strength is caused by both a restriction of dislocation movement due to HallPetch strengthening introduced by high strain rate deformation and a reduction in the dislocation mobility due to the interaction of dislocations with the nanoprecipitates introduced by DP. 4.4. Deeper plastically deformed layer from LA-UNSM In addition to the unique nanostructure, LA-UNSM produced a thicker SPD zone as compared with that of UNSM as shown in Figs. 8b and 8c. It is known that annealing normally results in material softening, leading to a decrease in hardness and a release of residual stress. However, an optimal working temperature can facilitate the plastic deformation by decreasing the flow stress [12,13]. During LA-UNSM, the saturated defects introduced by the ultrasonic strikes were partially annihilated due to the laser heating, which leaves room for further plastic deformation. The high-frequency strikes of the ultrasonic tip effectively superimposes new plastic deformation on the softened material. Meanwhile, due to the accelerated atom migration and dislocation movement, the plastic deformation propagates deep into the material and thereby giving rise to a thicker SPD layer as compared to that of the UNSM-treated Ti64. 4.5. Local heating Another advantage of laser heating is the ability to heat localized areas without affecting the entire workpiece. Elevated temperature annealing normally results in recovery, recrystallization, and grain coarsening. However, TEM images (Figs. 10 and 11) demonstrate that the grain size in the LA-UNSM-processed Ti64 sample is similar to that found in the UNSM-processed sample, indicating similar degree of refinement. This is because although the maximum temperature during LA-UNSM was higher than 450 °C, this high-temperature field was confined to the small tip/sample contact area. Meanwhile, the laser beam moved with the tip simultaneously. That is to say, after the laser beam scans over an area, the mass of the material is generally sufficient for 20

ACCEPTED MANUSCRIPT

rapid heat removal. The region far away from the tip/workpiece contact area then remained at a comparatively low temperature (100°C~300°C) with a mild heating effect throughout the treatment. That is why no obvious grain growth was found in the LA-UNSM processed Ti64. Therefore, LA-UNSM can take advantage of the decrease in flow stress and overcome the limit of work-hardening. It should be noted that an increase in the laser power causes a rise in the surface hardness; however, when a high laser power was used (i.e., 68 W), a decrease in hardness was observed (Fig. 2a). This occurs because the degree of the annihilation of dislocations becomes high even as the magnitude of the plastic deformation increases with the rise in working temperature [12,13]. The strengthening brought about by work hardening cannot compensate for the loss induced by laser annealing. Additionally, for LA-UNSM at 68 W, the peak intensity of the precipitate (as shown in Fig. 14b) is lower than that of the LA-UNSM at 56 W. Since all XRD characterizations were carried out for workpieces of the same size, it is reasonable to conclude that the volume fraction of the precipitate in the Ti64 sample processed by LA-UNSM at 68 W is lower than that of the one processed by LA-UNSM at 56 W. The volume fraction decrease of the precipitate directly led to the decrease in strength. Thus, for LA-UNSM, the optimum laser power is 56 W in this study. Conclusion In this study, the microstructure evolution in Ti64 subjected to simultaneous ultrasonic nanocrystal surface modification and laser heating was investigated. It was observed that a composite microstructure characterized by a nanocrystalline layer embedded with nanoscale precipitates was produced in Ti64 after LA-UNSM. This unique microstructure resulted in a 75% increase in hardness. The following conclusions can be made: 1. A unique nanocomposite microstructure, which features a nanocrystalline layer embedded with nanoscale precipitates, can only be achieved through simultaneous high strain rate plastic deformation and laser heating in LA-UNSM. This hybrid process resulted in a significant improvement in the surface strength that is much higher than that produced using traditional SSPD processes; 2. The nanoscale precipitates were produced through dynamic precipitation in laser-assisted high strain rate plastic deformation; 3. Both UNSM and LA-UNSM can refine the grains to the nanometer range. The local heating caused by the laser beam has a modest effect on grain growth; 4. Compared with UNSM, LA-UNSM substantially increases the plastic affected depth, since more plastic deformation can be introduced into the material due to the softening induced by laser heating.

21

ACCEPTED MANUSCRIPT

References [1]

J.Z. Lu, K.Y. Luo, Y.K. Zhang, G.F. Sun, Y.Y. Gu, J.Z. Zhou, et al., Grain refinement mechanism of multiple laser shock processing impacts on ANSI 304 stainless steel, Acta Mater. 58 (2010) 5354–5362. doi:10.1016/j.actamat.2010.06.010.

[2]

C. Ye, S. Suslov, X. Fei, G.J. Cheng, Bimodal nanocrystallization of NiTi shape memory alloy by laser shock peening and post-deformation annealing, Acta Mater. 59 (2011) 7219–7227. doi:10.1016/j.actamat.2011.07.070.

[3]

L. Zhou, C. Long, W. He, L. Tian, W. Jia, Improvement of high-temperature fatigue performance in the nickel-based alloy by LSP-induced surface nanocrystallization, J. Alloys Compd. 744 (2018) 156–164. doi:10.1016/j.jallcom.2018.01.070.

[4]

J. Lindemann, C. Buque, F. Appel, Effect of shot peening on fatigue performance of a lamellar titanium aluminide alloy, Acta Mater. 54 (2006) 1155–1164. doi:10.1016/j.actamat.2005.10.043.

[5]

D.J. Child, G.D. West, R.C. Thomson, Assessment of surface hardening effects from shot peening on a Ni-based alloy using electron backscatter diffraction techniques, Acta Mater. 59 (2011) 4825–4834. doi:https://doi.org/10.1016/j.actamat.2011.04.025.

[6]

X. Wu, N. Tao, Y. Hong, B. Xu, J. Lu, K. Lu, Microstructure and evolution of mechanically-induced ultrafine grain in surface layer of AL-alloy subjected to USSP, Acta Mater. 50 (2002) 2075–2084. doi:10.1016/S1359-6454(02)00051-4.

[7]

M.A. Vasylyev, S.P. Chenakin, L.F. Yatsenko, Ultrasonic impact treatment induced oxidation of Ti6Al4V alloy, Acta Mater. 103 (2016) 761–774. doi:10.1016/j.actamat.2015.10.041.

[8]

N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, An investigation of surface nanocrystallization mechanism in Fe induced by surface mechanical attrition treatment, Acta Mater. 50 (2002) 4603–4616. doi:10.1016/S1359-6454(02)00310-5.

[9]

K.Y. Zhu, A. Vassel, F. Brisset, K. Lu, J. Lu, Nanostructure formation mechanism of αtitanium using SMAT, Acta Mater. 52 (2004) 4101–4110. doi:10.1016/j.actamat.2004.05.023.

[10]

A. Amanov, I.S. Cho, Y.S. Pyoun, C.S. Lee, I.G. Park, Micro-dimpled surface by ultrasonic nanocrystal surface modification and its tribological effects, Wear. 286–287 (2012) 136–144. doi:10.1016/j.wear.2011.06.001.

[11]

A. Amanov, I.S. Cho, D.E. Kim, Y.S. Pyun, Fretting wear and friction reduction of CP titanium and Ti-6Al-4V alloy by ultrasonic nanocrystalline surface modification, Surf. Coatings Technol. 207 (2012) 135–142. doi:10.1016/j.surfcoat.2012.06.046.

[12]

Y. Harada, K. Mori, Effect of processing temperature on warm shot peening of spring steel, J. Mater. Process. Technol. 162–163 (2005) 498–503. 22

ACCEPTED MANUSCRIPT

doi:10.1016/j.jmatprotec.2005.02.095. [13]

G. Li, S. Qu, M.X. Xie, X. Li, Effect of ultrasonic surface rolling at low temperatures on surface layer microstructure and properties of HIP Ti-6Al-4V alloy, Surf. Coatings Technol. 316 (2017) 75–84. doi:10.1016/j.surfcoat.2017.01.099.

[14]

J.Z. Lu, H.F. Duan, K.Y. Luo, L.J. Wu, W.W. Deng, J. Cai, Tensile properties and surface nanocrystallization analyses of H62 brass subjected to room-temperature and warm laser shock peening, J. Alloys Compd. 698 (2017) 633–642. doi:10.1016/j.jallcom.2016.12.210.

[15]

C. Ye, Y. Liao, G.J. Cheng, Warm laser shock peening driven nanostructures and their effects on fatigue performance in Aluminum Alloy 6160, Adv. Eng. Mater. 12 (2010) 291–297. doi:10.1002/adem.200900290.

[16]

C. Ye, S. Suslov, B.J. Kim, E.A. Stach, G.J. Cheng, Fatigue performance improvement in AISI 4140 steel by dynamic strain aging and dynamic precipitation during warm laser shock peening, Acta Mater. 59 (2011) 1014–1025. doi:10.1016/j.actamat.2010.10.032.

[17]

Y. Liao, C. Ye, B.J. Kim, S. Suslov, E.A. Stach, G.J. Cheng, Nucleation of highly dense nanoscale precipitates based on warm laser shock peening, J. Appl. Phys. 108 (2010). doi:10.1063/1.3481858.

[18]

Y. Liao, C. Ye, H. Gao, B.J. Kim, S. Suslov, E.A. Stach, et al., Dislocation pinning effects induced by nano-precipitates during warm laser shock peening: Dislocation dynamic simulation and experiments, J. Appl. Phys. 110 (2011). doi:10.1063/1.3609072.

[19]

C. Ye, Y. Liao, S. Suslov, D. Lin, G.J. Cheng, Ultrahigh dense and gradient nanoprecipitates generated by warm laser shock peening for combination of high strength and ductility, Mater. Sci. Eng. A. 609 (2014) 195–203. doi:10.1016/j.msea.2014.05.003.

[20]

Y. Liao, S. Suslov, C. Ye, G.J. Cheng, The mechanisms of thermal engineered laser shock peening for enhanced fatigue performance, Acta Mater. 60 (2012) 4997–5009. doi:10.1016/j.actamat.2012.06.024.

[21]

H. Shipley, D. McDonnell, M. Culleton, R. Coull, R. Lupoi, G. O’Donnell, et al., Optimisation of process parameters to address fundamental challenges during selective laser melting of Ti-6Al-4V: A review, Int. J. Mach. Tools Manuf. 128 (2018) 1–20. doi:10.1016/j.ijmachtools.2018.01.003.

[22]

Z. Li, R. Xu, Z. Zhang, I. Kucukkoc, The influence of scan length on fabricating thinwalled components in selective laser melting, Int. J. Mach. Tools Manuf. 126 (2018) 1– 12. doi:10.1016/j.ijmachtools.2017.11.012.

[23]

H. Zhu, J. Wang, P. Yao, C. Huang, Heat transfer and material ablation in hybrid laserwaterjet microgrooving of single crystalline germanium, Int. J. Mach. Tools Manuf. 116 (2017) 25–39. doi:10.1016/j.ijmachtools.2017.01.002.

[24]

M. Masoomi, S.M. Thompson, N. Shamsaei, Laser powder bed fusion of Ti-6Al-4V parts: 23

ACCEPTED MANUSCRIPT

Thermal modeling and mechanical implications, Int. J. Mach. Tools Manuf. 118–119 (2017) 73–90. doi:10.1016/j.ijmachtools.2017.04.007. [25]

M. Anderson, R. Patwa, Y.C. Shin, Laser-assisted machining of Inconel 718 with an economic analysis, Int. J. Mach. Tools Manuf. 46 (2006) 1879–1891. doi:10.1016/j.ijmachtools.2005.11.005.

[26]

G. Germain, P. Dal Santo, J.L. Lebrun, Comprehension of chip formation in laser assisted machining, Int. J. Mach. Tools Manuf. 51 (2011) 230–238. doi:10.1016/j.ijmachtools.2010.11.006.

[27]

G.B.J. Cadot, D.A. Axinte, J. Billingham, Continuous trench, pulsed laser ablation for micro-machining applications, Int. J. Mach. Tools Manuf. 107 (2016) 8–20. doi:10.1016/j.ijmachtools.2016.04.011.

[28]

W. Cong, F. Ning, A fundamental investigation on ultrasonic vibration-assisted laser engineered net shaping of stainless steel, Int. J. Mach. Tools Manuf. 121 (2017) 61–69. doi:10.1016/j.ijmachtools.2017.04.008.

[29]

P. Liu, L. Deng, J. Duan, B. Wu, X. Zeng, Y. Shangguan, et al., A study on laser multifocus separation technology of thick KDP crystal, Int. J. Mach. Tools Manuf. 118–119 (2017) 26–36. doi:10.1016/j.ijmachtools.2017.04.002.

[30]

S.H. Masood, K. Armitage, M. Brandt, An experimental study of laser-assisted machining of hard-to-wear white cast iron, Int. J. Mach. Tools Manuf. 51 (2011) 450–456. doi:10.1016/j.ijmachtools.2011.02.001.

[31]

Y. Tian, Y.C. Shin, Laser-assisted burnishing of metals, Int. J. Mach. Tools Manuf. 47 (2007) 14–22. doi:10.1016/j.ijmachtools.2006.03.002.

[32]

A. Amanov, O. V. Penkov, Y.S. Pyun, D.E. Kim, Effects of ultrasonic nanocrystalline surface modification on the tribological properties of AZ91D magnesium alloy, Tribol. Int. 54 (2012) 106–113. doi:10.1016/j.triboint.2012.04.024.

[33]

M.A. Vasylyev, S.P. Chenakin, L.F. Yatsenko, Nitridation of Ti6Al4V alloy under ultrasonic impact treatment in liquid nitrogen, Acta Mater. 60 (2012) 6223–6233. doi:10.1016/j.actamat.2012.08.006.

[34]

S. Bahl, S. Suwas, T. Ungàr, K. Chatterjee, Elucidating microstructural evolution and strengthening mechanisms in nanocrystalline surface induced by surface mechanical attrition treatment of stainless steel, Acta Mater. 122 (2017) 138–151. doi:10.1016/j.actamat.2016.09.041.

[35]

Y. Liu, M. Li, Nanocrystallization mechanism of beta phase in Ti-6Al-4V subjected to severe plastic deformation, Mater. Sci. Eng. A. 669 (2016) 7–13. doi:10.1016/j.msea.2016.05.088.

[36]

C. Ye, A. Telang, A.S. Gill, S. Suslov, Y. Idell, K. Zweiacker, et al., Gradient 24

ACCEPTED MANUSCRIPT

nanostructure and residual stresses induced by Ultrasonic Nano-crystal Surface Modification in 304 austenitic stainless steel for high strength and high ductility, Mater. Sci. Eng. A. 613 (2014) 274–288. doi:10.1016/j.msea.2014.06.114. [37]

D. Ahn, Y. He, Z. Wan, I.S. Cho, C.S. Lee, I.G. Park, et al., Effect of ultrasonic nanocrystalline surface modification on the microstructural evolution and mechanical properties of Al5052 alloy, Surf. Interface Anal. 44 (2012) 1415–1417. doi:10.1002/sia.4959.

[38]

A. Amanov, Y.S. Pyun, S. Sasaki, Effects of ultrasonic nanocrystalline surface modification (UNSM) technique on the tribological behavior of sintered Cu-based alloy, Tribol. Int. 72 (2014) 187–197. doi:10.1016/j.triboint.2013.12.003.

[39]

J. Liu, S. Suslov, S. Li, H. Qin, Z. Ren, G.L. Doll, et al., Electrically Assisted Ultrasonic Nanocrystal Surface Modification of Ti6Al4V Alloy, Adv. Eng. Mater. (2017). doi:10.1002/adem.201700470.

[40]

J. Liu, S. Suslov, S. Li, H. Qin, Z. Ren, C. Ma, et al., Effects of ultrasonic nanocrystal surface modification on the thermal oxidation behavior of Ti6Al4V, Surf. Coatings Technol. 325 (2017) 289–298. doi:10.1016/j.surfcoat.2017.04.051.

[41]

H. Wang, G. Song, G. Tang, Evolution of surface mechanical properties and microstructure of Ti-6Al-4V alloy induced by electropulsing-assisted ultrasonic surface rolling process, J. Alloys Compd. 681 (2016) 146–156. doi:10.1016/j.jallcom.2016.04.067.

[42]

S. Mironov, M. Murzinova, S. Zherebtsov, G.A. Salishchev, S.L. Semiatin, Microstructure evolution during warm working of Ti-6Al-4V with a colony-α microstructure, Acta Mater. 57 (2009) 2470–2481. doi:10.1016/j.actamat.2009.02.016.

[43]

H. Shahmir, T.G. Langdon, Using heat treatments, high-pressure torsion and postdeformation annealing to optimize the properties of Ti-6Al-4V alloys, Acta Mater. 141 (2017) 419–426. doi:10.1016/j.actamat.2017.09.018.

[44]

F.X. Gil Mur, D. Rodríguez, J. a. Planell, Influence of tempering temperature and time on the α′-Ti-6Al-4V martensite, J. Alloys Compd. 234 (1996) 287–289. doi:10.1016/09258388(95)02057-8.

[45]

D. Lunt, T. Busolo, X. Xu, J. Quinta da Fonseca, M. Preuss, Effect of nanoscale α2 precipitation on strain localisation in a two-phase Ti-alloy, Acta Mater. 129 (2017) 72–82. doi:10.1016/j.actamat.2017.02.068.

[46]

D.G. Lee, S. Lee, Y. Lee, Effect of precipitates on damping capacity and mechanical properties of Ti-6Al-4V alloy, Mater. Sci. Eng. A. 486 (2008) 19–26. doi:10.1016/j.msea.2007.08.053.

[47]

W. Yu, M.Q. Li, J. Luo, Effect of deformation parameters on the precipitation mechanism of secondary α phase under high temperature isothermal compression of Ti-6Al-4V alloy, 25

ACCEPTED MANUSCRIPT

Mater. Sci. Eng. A. 527 (2010) 4210–4217. doi:10.1016/j.msea.2010.03.024. [48]

S.J. Lainé, K.M. Knowles, P.J. Doorbar, R.D. Cutts, D. Rugg, Microstructural characterisation of metallic shot peened and laser shock peened Ti–6Al–4V, Acta Mater. 123 (2017) 350–361. doi:10.1016/j.actamat.2016.10.044.

[49]

R.Z. Valiev, A.V. Korznikov, R.R. Mulyukov, Structure and properties of ultrafinegrained materials produced by severe plastic deformation, Mater. Sci. Eng. A. 168 (1993) 141–148. doi:10.1016/0921-5093(93)90717-S.

[50]

R.S. Mishra, V.V. Stolyarov, C. Echer, R.Z. Valiev, A.K. Mukherjee, Mechanical behavior and superplasticity of a severe plastic deformation processed nanocrystalline Ti– 6Al–4V alloy, Mater. Sci. Eng. A. 298 (2001) 44–50. doi:10.1016/S0921-5093(00)013381.

[51]

Y.G. Ko, W.S. Jung, D.H. Shin, C.S. Lee, Effects of temperature and initial microstructure on the equal channel angular pressing of Ti-6Al-4V alloy, Scr. Mater. 48 (2003) 197–202. doi:10.1016/S1359-6462(02)00356-1.

[52]

D.G. Lee, Y.G. Kim, D.H. Nam, S.M. Hur, S. Lee, Dynamic deformation behavior and ballistic performance of Ti-6Al-4V alloy containing fine α2 (Ti3Al) precipitates, Mater. Sci. Eng. A. 391 (2005) 221–234. doi:10.1016/j.msea.2004.08.076.

[53]

G. Welsch, G. Lütjering, K. Gazioglu, W. Bunk, Deformation characteristics of age hardened Ti-6Al-4V, Metall. Trans. A. 8 (1977) 169–177. doi:10.1007/BF02677278.

[54]

P.C. Sahu, N. V. Chandra Shekar, M. Yousuf, K. Govinda Rajan, Implications of a Pressure Induced Phase Transition in the Search for Cubic Ti3Al, Phys. Rev. Lett. 78 (1997) 1054–1057. doi:10.1103/PhysRevLett.78.1054.

[55]

E. Bonetti, G. Cocco, S. Enzo, G. Valdre, Amorphisation and phase transformations in mechanically alloyed Ti–Al powders: electron microscopy investigation, Mater. Sci. Technol. 6 (1990) 1258–1262. doi:10.1179/mst.1990.6.12.1258.

[56]

I. Manna, P.P. Chattopadhyay, P. Nandi, F. Banhart, H.J. Fecht, Formation of facecentered-cubic titanium by mechanical attrition, J. Appl. Phys. 93 (2003) 1520–1524. doi:10.1063/1.1530718.

[57]

M. Kheradmandfard, S.F. Kashani-Bozorg, C.L. Kim, A.Z. Hanzaki, Y.S. Pyoun, J.H. Kim, et al., Nanostructured β-type titanium alloy fabricated by ultrasonic nanocrystal surface modification, Ultrason. Sonochem. 39 (2017) 698–706. doi:10.1016/j.ultsonch.2017.03.061.

[58]

J.Z. Lu, L.J. Wu, G.F. Sun, K.Y. Luo, Y.K. Zhang, J. Cai, et al., Microstructural response and grain refinement mechanism of commercially pure titanium subjected to multiple laser shock peening impacts, Acta Mater. 127 (2017) 252–266. doi:10.1016/j.actamat.2017.01.050. 26

ACCEPTED MANUSCRIPT

[59]

Y.B. Chun, S.K. Hwang, Static recrystallization of warm-rolled pure Ti influenced by microstructural inhomogeneity, Acta Mater. 56 (2008) 369–379. doi:10.1016/j.actamat.2007.09.040.

[60]

Z. Zeng, Y. Zhang, S. Jonsson, Deformation behaviour of commercially pure titanium during simple hot compression, Mater. Des. 30 (2009) 3105–3111. doi:10.1016/j.matdes.2008.12.002.

[61]

S.Y. Mironov, G.A. Salishchev, M.M. Myshlyaev, R. Pippan, Evolution of misorientation distribution during warm “abc” forging of commercial-purity titanium, Mater. Sci. Eng. A. 418 (2006) 257–267. doi:10.1016/j.msea.2005.11.026.

[62]

B. Dutta, E. Valdes, C.M. Sellars, Mechanism and kinetics of strain induced precipitation of Nb(C,N) in austenite, Acta Metall. Mater. 40 (1992) 653–662. doi:10.1016/09567151(92)90006-Z.

[63]

E. Dogan, S. Wang, M.W. Vaughan, I. Karaman, Dynamic precipitation in Mg-3Al-1Zn alloy during different plastic deformation modes, Acta Mater. 116 (2016) 1–13. doi:10.1016/j.actamat.2016.06.011.

[64]

K.C. Russell, Nucleation in solids: The induction and steady state effects, Adv. Colloid Interface Sci. 13 (1980) 205–318. doi:10.1016/0001-8686(80)80003-0.

27

ACCEPTED MANUSCRIPT

Acknowledgements Authors J. Liu, Z. Ren, Y. Dong and C. Ye are grateful for the financial support of this research by the Start-up Fund and the Firestone Fellowship provided by the College of Engineering at The University of Akron. We also would like to acknowledge the Qatar Environment and Energy Research Institute for facility support of the TEM work.

ACCEPTED MANUSCRIPT

Highlights: 1. Laser assisted ultrasonic nanocrystal surface modification (LA-UNSM) integrates laser heating with UNSM; 2. In additional to nanocrystallization, LA-UNSM produce nanoscale precipitates through dynamic precipitation; 3. By fabricating a unique nanocomposite microstructure, LA-UNSM significantly increase the surface strength of Ti64; 4. As compared with UNSM, LA-UNSM can significantly increase the plastic affected depth due to material softening by laser heating; 5. Local heating from laser effectively suppress grain coarsening as compared with traditional furnace heating.