Microstructural and mechanical property evolution of a nuclear zirconium-4 alloy fabricated via laser powder bed fusion and annealing heat treatment

ABSTRACT Zirconium (Zr) alloys are widely used in nuclear energy because of their excellent mechanical properties and low thermal neutron absorption cross-section. This work investigated the printability, microstructure, and mechanical properties of Zr-4 alloy additively manufactured by laser powder bed fusion (LPBF) for the first time. The effect of annealing temperature on the microstructural and the mechanical property evolution of the printed Zr-4 alloy was studied. The results exhibited that the Zr-4 alloy with a high relative density of 99.77% was obtained using optimised printing parameters. With an increase in the annealing temperature, the formed α phase of the Zr-4 alloy changed from an acicular shape to a coarse-twisted shape, and finally to an equiaxed shape. Such microstructure change endowed the alloy with a high compressive strength of 2130 MPa and compressive strain of 36%. When the annealing temperature exceeded 700°C, Zr x (Fe2Cr) compounds were precipitated, strengthening the alloy by pinning effect. These findings provide valuable guidance for the manufacture of geometrically complex Zr alloy parts for nuclear power applications.


Introduction
Zirconium and its alloys are characterised by low thermal neutron absorption cross-section (0.18 barn), high hardness, ductility, and excellent corrosion resistance (Abe 1991).Generally, four types of microstructure can be found in Zr alloys, i.e. equiaxial structure, bimodal structure, basket-weave structure, and widmannstatten structure (Singh et al. 2013).The equiaxed structure is usually formed when the cooling temperature of Zr alloys is in the middle and low temperature range of the α+β region.The bimodal structure appears when the deformation temperature is in the high temperature range of the α+β region.The basket-weave structure can be obtained when the initial deformation temperature is slightly higher than β phase transition temperature, and the end temperature is in α+β range.The widmannstatten structure is obtained when the initial deformation temperature is much higher than the β phase transition temperature.
Zirconium-4 (Zr-4) alloys, which consist of Zr (>95 wt %), Sn, Nb, Fe, Cr, and Ni elements, serve as representative candidate materials in the nuclear industry (Linga Murty and Charit 2006).Zr-4 alloys have been applied for the manufacture of fuel cladding (acting as a longcycle and high-burnup fuel component), core structures of nuclear power reactors, nuclear spacer grids, element boxes, heat exchangers, and core structures of pressurised water reactors.A minimum of 400 tons of Zr-4 alloys is required for a million-scale nuclear power unit during its service (Zinkle and Was 2013;Shi et al. 2013).
Zr-4 alloys are commonly manufactured by conventional methods such as casting, forging, and rolling.Fuloria et al. evaluated the mechanical properties and microstructure evolution of multiaxial-forged Zr-4 alloys under different strains at low temperatures (Fuloria et al. 2015).With an increase in the strain, the ultimate tensile strength and hardness of the alloy increased from 474 to 717 MPa and from 190 HV to 238 HV, respectively.Fuloria et al. also investigated the effects of hot rolling temperatures and deformation strains on the hardness and tensile properties of Zr-4 alloys.The results revealed that an increase in the stacking fault probability at increased temperatures during the progressive deformation resulted in enhanced strength and hardness (Fuloria et al. 2016).Sarkar et al. studied the evolution of microstructure and mechanical properties of rolled Zr-4 alloys at low temperatures, and found that an increase in the dislocation density enhanced the strengthening effect of the alloy (Sarkar and Murty 2015).However, it is still challenging to fabricate Zr-4 nuclear parts with complex geometries and superior mechanical performance using these conventional methods.
Additive manufacturing (AM) is capable of fabricating geometrically complex parts in a layer-by-layer manner (Ngo et al. 2018;Wu, Su, and Shi 2019).AM can be used to fabricate multi-materials to satisfy the complex requirements of industries (Li et al. 2022a) and print alloys and composites through in-situ fabrication (Li et al. 2022b;Mosallanejad et al. 2021;Sing et al. 2021).Directed energy deposition (DED) and laser powder bed fusion (LPBF) are the popular metal AM techniques for fabricating nuclear parts (Zhao and Wang 2022;Smith et al. 2021).LPBF uses highenergy laser beams to melt metal powders (Liu et al. 2022;Xue et al. 2022).The small laser spot size and powder particle size in LPBF enable the high-precision printing of metal parts, while DED possesses faster build rate and larger build volume (DebRoy et al. 2018;Yu et al. 2022).Several studies have been reported on Zr alloys printed with the above two processes.For instance, Su et al. investigated the tensile properties of a Zr-based bulk metallic glass prepared by DED (Su and Lu 2019).They achieved a high ultimate tensile strength of 880 MPa and high yield strength of 835 MPa, which were comparable to those of the as-cast counterpart.
Compared to DED, LPBF possesses higher melting and solidification rates in a much smaller melt pool and allows a more precise energy control.Currently, there are only a few research reports on Zr-4 alloys fabricated by LPBF.Hence, in this work, the printability, microstructure, and mechanical properties of a Zr-4 alloy fabricated by LPBF were investigated for the first time.The LPBF process parameters for the Zr-4 alloy were optimised to evaluate its printability.The effect of annealing heat treatment on the microstructure and mechanical properties of the printed alloy was further studied.The underlying mechanisms of the microstructure evolution and property enhancement of the alloy were analysed and discussed.

Materials
Spherical Zr-4 alloy powder (supplied by Hebei Baoju New Material Technology Co., Ltd) prepared by plasma milling with a particle size distribution of 15-53 μm was used.The chemical composition of the powder is listed in Table 1 and the powder morphology is shown in Figure 1(a).A laser diffraction particle size analyzer (HORIBA LA960S) was used to measure the powder diameter distribution of Zr-4 alloys.The average particle diameter was 31.35 μm (Figure 1(b)).The flowability of the powder was evaluated by the Hall velocity value measured using a Hall flow meter (KC-02) (Zhang, Gu,  and Dai 2022; Yu et al. 2022).The flowability of the Zr-4 powder was measured to be 17.52 s/50 g, which meets the requirement of nuclear-graded powders (24 s/50 g) (Sehra, Vijay, and Gupta 1995).

LPBF process optimisation
A Dimetal-100H LPBF machine (Laseradd, Guangzhou, China) was employed to fabricate the Zr-4 alloy from its powder (Figure 2(a)).An orthogonal scanning strategy of 90°transition between layers was adopted (Cheng, Shrestha, and Chou 2016) (Figure 2(b)).The models and printed Zr samples for testing are shown in Figure 2(c,d), respectively.The LPBF chamber was filled with argon to ensure that the oxygen content was less than 0.01 vol%.Based on our previous exploration on process parameters (Table S1), a set of orthogonal experiments was designed for process parameter optimisation, as listed in Table 2.The layer thickness and spot diameter were set as 0.03 and 0.08 mm, respectively.The volumetric energy density E v (J/mm 3 ) can be described as: where P is the laser power (W), V is the scanning speed (mm/s), S is the hatch space (mm) and h is the layer thickness (mm).The experimental density of the samples was measured by the Archimedes drainage method, and their relative density was determined by the ratio of the measured density to the theoretical density (6.5 g/ cm 3 ).Three sets of data were measured from the samples to determine the average relative density value.

Annealing process
The annealing process was conducted in a furnace with vacuum environment (N41/H heat treatment furnace, Nabertherm).The Zr-4 samples printed with optimised process parameters were annealed at the temperatures of 580°C, 650°C, 700°C, 750°C, 800°C, 900°C, and 1000°C.The Zr-Sn phase diagram was referred to guide the selection of the annealing temperatures (Figure 3(a)).The annealing process under various temperatures is shown in Figure 3(b).A heating rate of 10°C /min and holding time of 2 h were applied to each temperature group, followed by air cooling.

Characterisations
LPBF-printed Zr-4 samples with and without annealing were prepared for metallographic observation.The samples with dimensions of 10 mm × 10 mm × 10 mm were ground, polished, and eventually etched by a corrosive solution (10 vol% HF, 45 vol% HNO 3, and 45 vol% H 2 O) for 15 s.The porosity was measured using an X-ray detector (Phoenix V|tome|x).An optical microscope (OM, DMI-3000M) and a scanning electron microscope (SEM, FEI Quanta 250) equipped with energy dispersive analyzer (EDS) were used to observe the microstructure of the samples.The element distributions of the samples were detected with an EPMA-1600 electron probe.The phase composition was identified through X-ray diffraction (XRD, X'Pert3 Powder diffractometer).A transmission electron microscope (TEM, TESCNAI G2F30, Thermo Fisher) was applied to further analyse the segregation of alloy elements and the crystal structure of the samples.Focused ion beam (FIB, Helios 600i) was employed to prepare TEM samples.
Electron backscatter diffraction (EBSD) was conducted using an SEM system (EDAX HIKARI SERIES) to characterise the grain characteristics of the samples.The samples were finely ground, polished with alumina liquid, perchloric acid, and methanol (20:80) etchant, and electropolished at 20 V for 1 min.A voltage of 20 kV and the minimum scanning step size of 0.1 μm were adopted.The analysis of the EBSD data was performed using a TSL OIM software.

Mechanical testing
Mechanical properties of the printed Zr-4 samples were tested using a universal mechanical machine (CMT5504) with a load capability of 50 kN under a constant stretch rate of 0.2 mm/min.Three compressive samples for each different annealing process were tested to estimate the compression properties according to the GB/T 7314-2017 standard.The hardness (HV) was tested with a load of 0.3 kg and a loading time of 10 s (WILSON VH1202 microhardness tester).Ten points were tested for each group.The compressed fractures were observed by SEM.

Process optimisation for the Zr-4 alloy
Figure 4(a) shows the variation of the relative density of the LPBF-printed Zr-4 alloy against the energy density.When the energy density range was from 133.3 to 166.7 J/mm 3 , the printed Zr-4 alloy with relative density greater than 99% was achieved.Figure 4(b) presents an iso-density diagram of the relative density as a function of laser power and scanning speed.The hatch space was fixed as 0.06 mm according to the orthogonal experiments results shown in Figure S1.The highest relative density was obtained under the laser power range of 190-230 W and the scanning speed of 600-700 mm/s.Figure 4(c) presents the OM images showing the pore morphology at different energy densities.When the energy density gradually increased to 106.7 J/mm 3 , some unmelted powder particles caused by insufficient energy densities decreased, contributing to increasing the relative density.However, as the energy density further increased to 186.7 J/mm 3 , the laser energy accumulated in the molten pool induced balling phenomenon and the formation of large-size pores.Thus, the relative density decreased (Sun et al. 2020a;Shelton et al. 2021;Carter et al. 2016).Figure 4(d) shows 3D reconstruction of the printed alloy with the highest relative density of 99.77% and the average pore diameter of 40.10 μm.Due to the process characteristics of rapid melting and rapid solidification, the protective gas cannot escape in time due to the viscosity of the liquid Zr-4 molten pool during LPBF process (Tan et al. 2021).Therefore, gas was trapped in the molten pool to form gas pores to decrease the density (Tan, Kiran, and Zhou 2021).Consequently, the optimised process parameters for the LPBF-printed Zr-4 alloy were determined as the laser power of 200 W, scanning speed of 650 mm/ s, hatch space of 0.06 mm, and layer thickness of 0.03 mm.
The LPBF-printed Zr-4 samples were subsequently heat-treated by different annealing process parameters.The variation curve of the relative density of the printed Zr-4 alloy at different annealing temperatures is presented in Figure 4(e).With the increase in annealing   temperature, the relative density of the alloy slowly increased and reached a peak value of 99.92% at 900°C , and then decreased as the temperature rose to 1000°C.Such density decrease may be attributed to the phase transformation in the Zr-4 alloy (Samanta and Kumar 2022).The effect of the annealing process on the relative density was determined by the healing of internal pores.With the increase in annealing temperature, the diffusion and dissolution of alloy elements were accelerated.Thus, internal pores were reduced by volume diffusion coupled with α-grain boundary diffusion (Zhang and Sun 2003).
Changes in process parameters lead to the variation of energy density.Low energy densities lead to unmelted powder particles due to insufficient energy input and balling phenomenon.In contrast, high energy densities may lead to gas pores and stress cracking resulting from the large residual stress caused by the large temperature gradients (Kuo et al. 2019;Guo et al. 2020).These defects negatively impact the relative density and mechanical properties of Zr-4 alloy, and they cannot be eliminated by annealing.Therefore, it is necessary to conduct parameter optimisation before the annealing treatment.

Phase identification
Figure 5 exhibits the XRD patterns of the Zr-4 alloy powder samples, and heat-treated samples under different annealing temperatures.The results indicated that the dominant phase constituent of Zr-4 alloy was α-Zr phase.The allotropic transformation of α-Zr to β-Zr can occur to generate β-Zr at temperatures higher than the phase transition temperature of Zr alloy.However, no β-Zr peaks were identified in the XRD patterns of the Zr-4 alloy annealed at 900°C and 1000°C.Since β-Zr is a stable phase at high temperatures, the phase transition from β to α easily occurred during air cooling (Devi et al. 2020), accompanied by non-equilibrium transformation and formation of metastable structures (such as martensite).The phase transition temperature of the Zr-4 alloy was increased by addition of Cr (α-Zr stabilising element) while it decreased by addition of Fe (β-Zr stabilising element).A small amount of Fe may result in generating α'/α'' martensite phase (similar crystal parameters as α phase) during air cooling.Therefore, the microstructure at 900°C and 1000°C remained the α phase, as reported in previous studies (Rogachev et al. 2017;Zhang et al. 2017;Feng et al. 2016).
Figure 6 presents the TEM images showing the phases and their crystal orientations in the LPBFprinted Zr-4 alloy processed with and without annealing at 700°C.The bright field images of the printed alloy and its annealed sample are shown in Figure 6(a-b), respectively.The as-printed Zr-4 alloy contained the α-Zr phase

Microstructure
Figure 7 displays the OM images showing morphology evolution of the LPBF-printed Zr-4 alloy with optimised parameters, processed with and without annealing.The typical morphology of the printed Zr-4 alloy (Figure 7(a,b)) showed good overlapping between the melt tracks and the matrix of the α phase.After annealing at 580°C, the melt tracks in both transverse and longitudinal planes partially disappeared (Figure 7(c)).When the annealing temperature increased to 650 and 700°C, the α phase became twisted and coarser (Figure 7(d,e)).Further increase in annealing temperature to 750°C resulted in the recrystallisation of the α phase and formation of sporadic equiaxed morphology.Most equiaxed α-grains appeared at the boundary of the α phase (Figure 7(f)).The recrystallisation was accelerated at the temperature of 800°C, and the volume fraction of the equiaxed α phase was significantly increased (Figure 7(g)).At the annealing temperature of 900°C, which exceeded the β phase transition temperature, the braided-shaped α phase was completely transformed into the equiaxed α phase (Figure 7(h)).When the temperature further increased to 1000°C, the equiaxed α grains with refined sizes were distributed homogeneously (Figure 7(i)).
Figure 8 presents the SEM images of the microstructure of the LPBF-printed Zr-4 alloy processed with and without annealing.As shown in Figure 8(a), the microstructure of the printed Zr-4 alloy was dominated by the acicular α phase.The α-grain boundaries of the printed alloy appeared at annealing temperatures exceeding 650°C (Figure 8(bd)).Additionally, an increase in the annealing temperature tended to promote microstructure homogenisation (Sun et al. 2017) and the precipitation of second phases at the αgrain boundaries.Above 800°C, a large amount of precipitates were located at the α-grain boundaries (Figure 8(e,f)).EDS analysis of the sample annealed at 900°C (Figure 8(g)) indicated that the content of Fe and Cr in the precipitated phase was greater than that in the matrix (Figure 8(h,i)), which can be presumed to be Zr x (Fe 2 Cr) compounds (Tao et al. 2018;Chen et al. 2018;Liu et al. 2019Liu et al. , 2018)).
Figure 9 presents the EBSD results of the LPBF-printed Zr-4 alloy processed without and with annealing at 700°C . As seen from Figure 9(a,b), the acicular α-grain had an average size of 1.51 μm in the printed Zr-4 alloy, which was consistent with the SEM result (Figure 8(a)).This αgrain exhibited a maximum pole density of 6.834 (Figure 9(c)).Comparatively, the α-grain of the annealed Zr-4 ally began to recrystallise and tended to be equiaxed (Figure 9(d)), which was consistent with the observations in Figures 7 and 8. Significantly, the average α-grain size of 31.4 μm was obtained (Figure 9 (e)), indicating the α-grain growth at 700°C.In addition, the grain orientation of the annealed alloy was random, with a maximum pole density of 6.08 (Figure 9(f)).In the printed Zr-4 alloy, the acicular α-grain nucleated along the direction of temperature gradient.However, when annealed at 700°C, the acicular α-grain of the printed alloy became coarser and began to recrystallise.This resulted in a significant increase in the α-grain size and the disordered grain orientations (Sun et al. 2021;Zhu et al. 2018;Hu et al. 2020).
The distribution of Sn and Fe elements in the printed Zr-4 alloy at the annealing temperatures of 650°C, 700°C, and 800°C is shown in Figure 10.The samples annealed at 650°C, 700°C, and 800°C were detected to analyse the distribution of alloy elements, as shown in Figure 10(ac), respectively.Fe element was concentrated at the α-grain boundary of the printed Zr-4 alloy with the increase in annealing temperature (Figure 10(df)).Sn-enriched areas appeared at the temperatures of 650°C and 700°C but disappeared at 800°C (Figure 10 (gi)).
Figure 11 exhibits the TEM images showing the elemental analysis of the printed Zr-4 alloy processed with and without annealing.Angle dark field (ADF) images were acquired with a DF4 probe in scanning transmission electron microscope (STEM) mode.It was observed that the α-grain size of the annealed Zr-4 alloy was much larger than that of its counterpart without annealing (Figure 11(a,b)).Moreover, Zr, Fe, Cr, and Sn elements were uniformly distributed in the printed Zr-4 matrix (Figure 11(c)).After annealing, the α-Zr martensite phase remained in the Zr-4 alloy, while Fe and Cr elements were segregated to form nanoscale Fe-rich and Cr-rich precipitates (Figure 11(d)).As seen from the EDS analysis results, the Fe and Cr contents in the printed alloy increased from 0.07 to 0.29 wt% and from 0.03 to 0.08 wt% after annealing, respectively (Figure 11(eg)).
Annealing temperature had a significant effect on the microstructure of the LPBF-printed Zr-4 alloy, as depicted in Figure 12.When the annealing temperature ranged from 580°C to 700°C, the dendritic  microstructural features of the printed Zr-4 alloy disappeared, and coarser α martensite phases were obtained (Figures 7 and 8).After the annealing temperature reached 750°C, the α phase began to recrystallise, and the equiaxed α phase appeared sporadically.Further increase in the annealing temperature promoted the recrystallisation and the precipitation of the second phases at the α-grain boundaries (Figure 8(d)).When the annealing temperature exceeded 900°C, the braided α phases completely evolved into the equiaxed morphology, illustrating that the recrystallisation process was completed (Figure 7(h) and Figure 8(f)).When the temperature was further increased to 1000°C , the sizes of the equiaxed α phase were refined (Figure 7(i)), and a large number of Zr-based compound precipitates appeared at both the boundaries and interior of the α-grain.Such grain refinement can be attributed to the pinning effect of the precipitates (Sun et al. 2020b(Sun et al. , 2020c(Sun et al. , 2021;;Jiang et al. 2018;Li et al. 2022a).

Mechanical properties
Figure 13 presents the Vickers hardness results of the LPBF-printed Zr-4 alloy processed with and without annealing.The printed alloy possessed the highest hardness of 397 HV.When the annealing temperature increased to 700°C, the hardness of the alloy gradually decreased to 348 HV.However, a further increase in the annealing temperature to 1000°C resulted in hardness improvement to 375 HV.The highest hardness of the printed Zr-4 alloy was attributed to the acicular martensite grains with small sizes and a large amount of residual stress existing in the alloy (Li et al. 2021a).During the annealing process, below 700°C, the residual stress was gradually released and the coarsening of α-grain occurred, resulting in lower hardness.Annealing temperatures above 700°C enhanced the atomic diffusion ability and the precipitation, promoting the dispersion strengthening effect.In addition, when the annealing temperature increased to 1000°C, the refinement of the equiaxed α-grain from full recrystallisation led to the increase in hardness.
Figure 14 shows the compressive properties of the LPBF-printed Zr-4 alloy with and without annealing.The compression data is shown in Table 3.As shown in Figure 14(a, b), the printed Zr-4 alloy possessed a compressive strength of 1650 MPa and a compressive strain of 19%.When the annealing temperature ranged from 580°C to 700°C, the compressive strength and compressive strain of the alloy both increased from 1710 to 2130 MPa and from 19% to 36%, respectively.The printed Zr-4 alloy annealed at 700°C achieved the highest strength and ductility.Above 700°C, the compressive strength of the annealed alloy decreased with the increase in annealing temperature.However, the compressive strain remained consistent from 700°C to 900°C, but reduced to 34% at 1000°C.
The printed alloy exhibited both low compressive strength and strain due to a large amount of residual  stress and relatively consistent grain orientation.Cracks were induced by stress concentration, which grew rapidly between grains with similar orientation and led to poor compressive properties of the printed alloy.When the annealing temperature increased to 700°C, the α-grain became larger and therefore hampered the migration of the dislocation, helping to improve the compressive strength of the alloy.Additionally, with the increase in the annealing temperature, the increasing lamellar α'/α'' phase produced more boundaries between the original α phase and α'/α'' phase, resulting in great resistance for movement of dislocations (Wang et al. 2018).Additionally, the temperature increase released the residual stress existing in the printed alloy, which was beneficial to ductility enhancement (He, Chen, and Liu 2018).When the annealing temperature exceeded 700°C, the αgrain began to recrystallise and the precipitates were increased.The increased precipitates had an intensified pinning effect on the alloy, resulting in compressive strength larger than 2 GPa.However, the increase in annealing temperature above 700°C also facilitated the recrystallisation and coarsened the α-grain (Shen et al. 2021;Li et al. 2021bLi et al. , 2022b)), decreasing the mechanical strength.
Figure 15 displays the representative fracture morphology of the LPBF-printed Zr-4 alloy with and without annealing at 700°C after compressive testing.A 45°shear fracture surface was observed in the printed and annealed alloy (Figure 15(a, d)).For the printed alloy, quasi-cleavage (Figure 15(b)) and cleavage surfaces (Figure 15(c)) appeared alternately, which was consistent with its low compressive strain.The alloy annealed at 700°C possessed more quasi-cleavage surfaces (Figure 15(e, f)), contributing to its high compressive strength (Ma et al. 2021).

Conclusions
This work investigated the printability, microstructure, and mechanical properties of a Zr-4 alloy additively manufactured by LPBF followed by annealing heat treatment.The effect of annealing temperature on the microstructure and properties of the printed alloy was evaluated.The main findings are presented as follows: (1) The LPBF-printed Zr-4 alloy exhibited a high relative density of 99.77% using optimised parameters, i.e. laser power of 200 W, scanning speed of 650 mm/ s, hatch space of 0.06 mm, and powder layer thickness of 0.03 mm.Energy density lower than 106.7 J/mm 3 led to low relative densities due to the existence of unmelted Zr-4 powder particles.In contrast, energy density higher than 186.7 J/mm 3 caused the balling phenomenon which decreased the relative density.
(2) The microstructure of the printed Zr-4 alloy was dominated by the α phase.The microstructure evolution of the printed alloy was determined by the annealing process.Increase in the annealing temperature enabled the microstructure evolution from the acicular α-phase to coarse and twisted α-phase as well as to equiaxed α-phase.This was accompanied by the increase in the content of precipitates that were likely Zr x (Fe 2 Cr) compounds.
(3) As the annealing temperature increased to 700°C, the printed Zr-4 alloy showed the highest compressive strength of 2130 MPa and compressive strain of 36%.These excellent mechanical properties may be attributed to grain boundary strengthening and precipitation strengthening.However, the increase in annealing temperature from 700°C to 1000°C also facilitated the recrystallisation, decreasing the mechanical strength.The results may be valuable for the development of Zr-4 alloy for nuclear power applications.

Disclosure statement
No potential conflict of interest was reported by the author(s).

Funding
This

Figure 2 .
Figure 2. (a) Dimetal-100 SLM equipment, (b) orthogonal scanning strategy, (c) geometries of the compressive sample model; (d) image of the printed Zr-4 tube parts and samples for compression testing.

Figure 4 .
Figure 4. Printability of Zr-4 alloy prepared by LPBF: (a) Variation of relative density against energy density; (b) iso-density diagram (hatch space of 0.06 mm); (c) OM images showing pore morphology of the alloy with different energy densities; (d) 3D reconstruction obtained from Zr-4 alloy printed with optimised parameters from Micro-CT; (e) relative density under different annealing temperatures.

Figure 5 .
Figure 5. XRD patterns of the Zr-4 alloy samples with different annealing temperatures.

Figure 6 .
Figure 6.TEM images showing the crystal orientation of the LPBF-printed Zr-4 alloy: (a, b) bright field image of the alloy processed without or with annealing at 700°C.Selected area electron diffraction (SAED) images of (c, d) regions A and B in image (a), respectively; (e-g) regions C, D and E in image b, respectively.

Figure 9 .
Figure 9. EBSD results of the LPBF-printed Zr-4 alloy processed with or without annealing.(a-c) IPF diagram, grain size distribution, and pole figure of the printed alloy, respectively.(d-e) IPF diagram, grain size distribution, and pole figure of the printed alloy annealed at 700°C, respectively.

Figure 10 .
Figure 10.EPMA analysis showing the element distribution of the printed Zr-4 alloy under different annealing temperatures: (a-c) the detection regions from samples with 650°C, 700°C, and 800°C annealing, respectively; (d-f) the distribution of Fe at the temperature of 650°C, 700°C, and 800°C, respectively; (gi) the distribution of Sn at the temperature of 650°C, 700°C, and 800°C, respectively.

Figure 11 .
Figure 11.TEM images showing the morphology and the element distribution of the LPBF-printed Zr-4 alloy with or without annealing at 700°C: (a, b) ADF morphology of the printed alloy and annealed alloy, respectively; (c, d) element distributions in the printed alloy and annealed alloy through EDS mapping, respectively; (e, f) element identification of the printed alloy and annealed alloy through EDS map analysis, respectively.

Figure 12 .
Figure 12.Schematic illustration of the microstructure evolution of the LPBF-printed Zr-4 alloy during annealing.

Figure 13 .
Figure 13.Vickers hardness changes of the LPBF-printed Zr-4 alloy with different annealing temperatures.

Figure 14 .
Figure 14.Compressive stress-strain responses of the LPBF-printed Zr-4 alloy without and with different annealing temperatures.

Table 3 .
Compressive properties of Zr-4 alloy with or without annealing.
work was supported by Guangdong Basic and Applied Basic Research Foundation (No. 2020B1515120013), Guangdong Province Science and Technology Project (No.