Microstructure and properties of 24CrNiMoY alloy steel prepared by direct laser deposited under different preheating temperatures

Journal Pre-proof Microstructure and properties of 24CrNiMoY alloy steel prepared by direct laser deposited under different preheating temperatures

Lin Zhou, Suiyuan Chen, Mingwei Wei, Jing Liang, Changsheng Liu, Mei Wang PII:

S1044-5803(19)31417-2

DOI:

https://doi.org/10.1016/j.matchar.2019.109931

Reference:

MTL 109931

To appear in:

Materials Characterization

Received date:

28 May 2019

Revised date:

10 September 2019

Accepted date:

11 September 2019

Please cite this article as: L. Zhou, S. Chen, M. Wei, et al., Microstructure and properties of 24CrNiMoY alloy steel prepared by direct laser deposited under different preheating temperatures, Materials Characterization (2019), https://doi.org/10.1016/ j.matchar.2019.109931

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© 2019 Published by Elsevier.

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Microstructure and Properties of 24CrNiMoY Alloy Steel Prepared by Direct Laser Deposited under Different Preheating Temperatures Lin Zhou 1, Suiyuan Chen1*, Mingwei Wei1, Jing Liang1, Changsheng Liu1, Mei Wang2 1-Key Laboratory for Anisotropy and Texture of Materials, Ministry of Education, Key Laboratory for Laser Application Technology and Equipment of Liaoning Province, School of Materials and Engineering, Northeastern University, Shenyang 110819, Liaoning, China.

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Email:[email protected].

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2-Shenyang Dalu Laser Technology Co. Ltd..

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Abstract

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*Corresponding author: Suiyuan Chen, email address: [email protected]

In order to effectively eliminate the crack initiation, reduce the amount of pores

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and improve the structure of large alloy steel parts with high-performance, the direct laser deposited (DLD) technique was used to prepare the 24CrNiMoY alloy steel

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samples by laser depositing power of 2000W and scanning speed of 6mm/s at 100℃,150℃, 200℃, 250℃ preheating temperatures, respectively. The structure and properties of the samples were studied using OM, SEM, XRD, TEM, EBSD, hardness tester and tensile tester. The results showed that with the increase of preheating temperature, the cracks in deposited alloy steel samples were removed, and the amount of pore was significantly reduced. The structure of the samples were mainly composed of granular bainite (GB) and lath bainite (LB), the content of LB was reduced from 95.7% to 3.5% with the increase of preheating temperature because GB played a segmentation role to LB. When the preheating temperature was 200℃, the deposited sample with GB (34.7%) and LB (65.3%) had the most uniform microhardness distribution, the highest microhardness of 414 HV0.2, and the best match of strength and toughness with the tensile strength of 1051 MPa and the

Journal Pre-proof elongation of 7.4%. The proper preheating temperature of 200℃ can effectively eliminate the crack, decrease the pores and significantly improve the structure and properties of the DLD 24CrNiMoY alloy steel.

Key words: Direct Laser Deposited; 24CrNiMoY alloy steel; Preheating; Bainite;

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Mechanical properties

Journal Pre-proof 1. Introduction In recent years, with the development of short-flow manufacturing technology, direct laser deposited (DLD) technology [1-3] has rapidly developed in rapid manufacturing [4-6] and rapid processing. DLD technology is the combination of the "stacking and accumulating" principle of rapid prototyping technology and laser cladding technology, raw materials for metal powder forming, with high energy laser beam as a heat source. According to the processing of forming parts computer aided design (CAD) model is stratified slice information path and send synchronization to

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metal powder for forming rapid solidification, depositing step by step, and realize the

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direct manufacturing of metal parts eventually. Compared with traditional casting technology, it has the advantages of saving materials, making complex parts and

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obtaining non-equilibrium solidified structure with fine grains and uniform

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composition [7]. However, DLD is a process of rapid heating melting, cooling and

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solidification, therefore, in the solidification process, the temperature gradient between the substrate and the deposited alloy layers is often too large, causing

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deformation, cracking, and generation of defects such as pores and inclusions. The predominant high heating and cooling rates and the existence of residual

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stress during deposition are the main reasons of crack initiation in Direct Laser Deposition (DLD) [8]. Cracks existed between the deposited layer and the substrate will lead to a significant drop in the mechanical properties of the deposited alloy part. Therefore, many researchers have tried various methods to prevent cracking emergence at the interface of the deposited layer, among which preheating substrate is one of the methods to reduce crack sensitivity [9-12]. The preheating substrate can reduce the temperature gradient between the substrate and the deposited layer during the deposition process, it can also decrease the thermal stress and residual stress by lowering the cooling rate, which is beneficial to slow the crack initiation. However, preheating also has an effect on the structure and properties of the deposited layer. Therefore, the preheating process and temperature should be reasonably selected to achieve the desired performance [13]. Fallsh et al. [14] studied the macroscopic and microscopic characteristics of

Journal Pre-proof Stellite 1 alloy coatings on preheated and non-preheated substrates, their experiments confirmed that preheating helped prevent cracking and provided a good profile on the surface of the sample, resulting in a more uniform surface hardness and dendritic distribution in the deposited sample. Shim et al. [15] preheated the substrate with an induction heater to study the effect of deposition characteristics on the high temperature tool steel M4 powder preheating, their experimental results showed that with the help of induction heating, preheating through the substrate could cause a slower cooling rate, which resulted in a coarser microstructure and a significantly

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longer spacing between the cells and the dendrite arms resulting in high

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microhardness. Liu et al. [16] investigated the effect of substrate preset temperature on crystal growth and microstructure formation in single crystal superalloy laser powder

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deposition, their results illustrated that the change of the substrate preset temperature

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between -30℃ and + 210℃ made the shape of the molten pool change little, but

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obviously affected the column-to-interchange conditions. Baek et al. [17] discussed the mechanical properties of high-speed cold stamping dies at different preheating

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temperatures, the upshot indicated that the thermal stress between the deposited layers was eliminated by preheating, and internal cracks were prevented. At the same time,

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the interface between the deposited layer and the substrate formed a dendritic structure, and a cell-like microstructure was formed in the deposited layer. As the preheating temperature increases, the cell crystals and precipitated carbides increased, and the hardness of the deposited layer slightly increased while the strength and toughness decreased.

Based on the previous studies, the research on the structure and properties of DLD alloy samples by using preheating method has achieved important results, which proves that the preheating method is an effective technique for controlling the crack defects of the deposited alloy samples and for the matching of structure and properties. However, the effect mechanism of preheating on crack elimination, pore formation and development, and microstructure and properties of DLD alloys is still insufficient. Especially, the effect of preheating treatment on the crack, structure and properties of laser-deposited 24CrNiMoY alloy steel samples has not been reported.

Journal Pre-proof Therefore, on the basis of our previous work [18-22], this paper adopts the preheating treatment method to study the amount of pore and crack, structure and property evolution of 24CrNiMoY alloy steel deposited by DLD. By designing different preheating temperatures and optimizing laser deposition process parameters, the best preheating treatment temperatures were obtained with usage of OM, XRD, SEM and other analytical methods. The internal defect content of deposited samples, bainite morphology, grain size, hardness and tensile properties were systematically investigated. The theory of crack elimination and change in pores content were

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studied, meanwhile the relationship between the preheating parameter and the change

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of structure and properties were established. The samples of 24CrNiMoY alloy steel with no defects and excellent structure performance were gained, which provides a

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reference for manufacturing large structural parts by DLD.

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2. Experimental methods 2.1 Materials

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24CrNiMoY alloy steel powder with particle size of 55-150 μm was used for the DLD process, the chemical composition of 24CrNiMoY alloy steel powder is

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shown in Table 1. The flowability of the powder was 15.1 s/50g, which was measured by the time when 50 g of powder was completely discharged from the Hall flow meter. The bulk density was 4.53 g/cm3, which was measured by freely filling the standard container with the powder under specified conditions. The powder was prepared by vacuum induction melting gas atomization (VIGA) method, which melted by laser heat source and solidified on the 100×200×10 mm3 size Q235 substrate. Table 1 Chemical compositions of as-used 24CrNiMo powder (wt.%) C

Cr

Ni

Si

Mn

Mo

Fe

Y

0.217

1.072

0.991

0.442

1.064

0.501

94.715

0.5

2.2 Experimental process Fig. 1 shows the DLD process design for preheating the substrate, using FL-Dlight02-3000W semiconductor laser (maximum power 3000 W, spot size 4 mm

Journal Pre-proof ×4 mm) and induction heater (the highest temperature is 400℃) to prepare samples. The induction heater was used for preheating the substrate in this study. The main heat transfer method is heat conduction, which is derived from the atomic activity in the form of lattice vibration. The substrate was preheated at the substrate temperature of 100℃, 150℃, 200℃ and 250℃ respectively. Firstly, the temperature was set by the control panel of the preheating device, and the substrate temperature was kept constant at the preset temperature (the error does not exceed 2℃). Then the deposition was processed with the continuous scan path by a semiconductor laser. The samples

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of the required size were cut and finally subjected to structure characterization and

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property testing.

Fig. 1. The DLD process diagram

2.3 Microstructural characterization and mechanical test After DLD manufacturing, the samples were cut perpendicular to the laser

Journal Pre-proof scanning direction (characterizing the YOZ plane) and their structure was observed. The preparation of the metallographic sample consisted of meshing the sample with different types of sandpaper, mechanical polishing with 2.5 micron diamond, then alcohol cleaning and drying, finally etching the sample with 4% nitric acid (4 ml nitric acid and 96 ml ethanol solution) for 10-15 s. The samples were characterized by OLYMPUS-GX71 optical microscope (OM), JSM-7001F field emission scanning electron microscope (SEM) with energy dispersive spectroscopy (EDS) attached. The parameters of EDS were 15 kV scanning voltage, 10 A current vacuum and 5.1×10-4

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vacuum degree. Then the image characteristics of the samples, such as single channel

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thickness, were analyzed and measured using Image Pro software. The EBSD sample was treated by electropolishing with 12.5% perchloric acid

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electrolyte (70 ml ethanol and 10 ml perchloric acid), and then subjected to electron

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backscatter diffraction using JSM-7001F with the scan step size of 0.2. X-ray diffraction was performed and analyzed using the following parameters: Cu target

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X-ray (λ = 0.154 nm), continuous scanning mode with the scanning angel ranged from 30° to 120°, sample scanning speed was 8 °/min, tube current and tube voltage

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were 40 mA and 40 kV, respectively. The TEM sample was first meshed to 30-50 μm

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by sandpaper, and then was punched into a thin disc with a diameter of 3 mm by a punching machine, after which it was processed by the 8% double spray (8% perchloric acid and 92% absolute ethanol) treatment with voltage 25 V, experiment temperature was -25°C and transmittance was 35. Density of alloy steel samples were obtained by Archimedes Drainage Test Method, the density of the sample can be obtained according to Eq. (1):



m1 m2  m3

(1)

where m1 is dry weight of sample tested with electronic balance instrument, m3 is floating weight in deionized water after vacuuming with a vacuum pump, m2 is wet weight by drying the surface of the sample. The Vickers microhardness of the cross section from surface to substrate was measured by a MicroMet-510 microhardness tester under the parameters of a 2 N load

Journal Pre-proof and a 10 s load-dwell time. The three sets of hardness data of each sample were input to the Origin software, the line tool was selected, and then the relevant parameter values of X and Y were set. The tensile sample was cut along the XOY plane (parallel to the scanning direction), and the geometry is shown in Fig. 2. The tensile test was carried out at room temperature using AG-X100KN at 0.5 mm/min. In order to ensure the accuracy of the experiment, three repeated tests were carried out under the same test conditions, and then a typical tensile stress-strain curve was obtained. The fracture of the sample was characterized by a JSM-6510A tungsten filament scanning

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electron microscope.

Fig. 2. Schematic of samples taken from deposition.

3. Results 3.1 Deposition characteristics of DLD 24CrNiMoY alloy steel samples 3.1.1 Cross-section of the deposited samples Large alloy steel parts tend to have crack initiation and some porosity defects during direct laser deposited process, which could significantly reduce the

Journal Pre-proof performance and life of the components. Fig. 3 shows the metallographic image of single layer of 24CrNiMoY alloy steel. As the rise of preheating temperature, the thickness of the single-layer deposition layer increased from 627 μm to 871 μm, and it can be seen from Fig. 3(e) that the single channel molten pool had a significant dilution at 250℃. At the same time, with the increase of preheating temperature, the amount of pores in the single-layer deposit first decreased obviously, and then increased at 250℃. In addition, micro-cracks appeared on one side of the melting pool of single deposition layer at Fig. 3(a), and the cracks disappeared after preheating the

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substrate, as shown in Fig. 3(b), (c), (d), (e). This phenomenon is attributed to the

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residual stress level remained in the samples when the temperature of the deposited samples was gradually lowered to the initial temperature. The specific reason is that

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the high-energy laser beam interacted with the deposition metal powder to form a

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liquid molten pool in the forming region, which was subjected to thermal expansion

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and restricted by the surrounding materials. When the laser beam was far away from the molten pool, the molten pool would instantly cool down, solidify and shrink in the

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volume, being constrained by the surrounding area on the deposited samples. These constraints in the deposition process led to complex internal stresses in the samples.

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When the complex internal stress is greater than the yield strength of the material, cracks of the deposited layer occurs [23,24].

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Fig. 3. Metallographic diagram of single channel deposition layer of 24CrNiMoY alloy steel: (a)Non-preheating, (b)100℃, (c)150℃, (d)200℃, (e)250℃.

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3.1.2 Formability of deposited samples

Fig. 4 is the macroscopic metallographic diagram of multilayer samples under

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different preheating parameters, it can be seen that there were pores remaining in the samples. The density of the samples shown in Fig. 5 is obtained according to Eq. (1), and the trend of it under different preheating temperature can be observed. It can be seen that the sample has the best compactness when the preheating temperature is 200℃. Besides, the distribution of the pores was not all at the bottom of the deposited layer, there were also pores in the middle and upper portions of the deposited layer, which confirms the floating of the pores in the samples during the DLD process. This phenomenon is mainly attributed to the generation of bubbles and molten pool movement during DLD process under different preheating treatment.

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Fig. 4. Macroscopic metallographic diagram of multi-layer deposits of 24CrNiMoY alloy steel: (a)Non-preheating, (b)100℃, (c)150℃, (d)200℃, (e)250℃.

Fig. 5. Density of DLD 24CrNiMoY alloy steel samples at different preheating temperature.

3.2 Structure evolution The X-ray diffraction phase analysis of DLD samples under different preheating temperature is shown in Fig. 6. It can be seen that the corresponding crystal plane indices of the three strong peaks were (110), (211), (310), the phase was mainly ɑ-Fe (M). Since both ferrite and bainitic ferrite (BF) are body-centered cubics structure (bcc), these strong peaks belong to the peaks of the two phases. Table 2 shows the crystallographic plane angles of 24CrNiMoY alloy steel sample.

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Fig. 6. X-ray diffraction analysis of 24CrNiMoY alloy steel with different preheating temperature

Table 2 Crystallographic plane angles of 24CrNiMoY alloy steel sample

100℃

150℃

200℃

250℃

44.52°

44.16°

44.16°

44.22°

64.64°

64.64°

64.54°

64.48°

64.46°

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25℃

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T/℃

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(HKL)

44.2°

(211)

82.06°

81.8°

81.86°

81.88°

81.78°

(220)

98.7°

98.36°

98.3°

98.28°

98.28°

(310)

116.08°

115.92°

115.9°

115.84°

115.8°

FWHM(110)

0.368°

0.303°

0.301°

0.301°

0.298°

(110)

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(200)

As shown in Table 2, the crystallographic plane angles of 24CrNiMoY alloy steel at different preheating temperatures showed that the peak value of XRD was offset to the left, which means the diffraction angel decreased. According to the Bragg equation, the crystal plane spacing became larger. This is because when the substrate was preheated, the stress level between the deposited layers during solidification was affected, which resulted in some macroscopic residual stress causing lattice distortion.

Journal Pre-proof In addition, FWHM of the samples decreased with the rise of preheating temperature. It can be known from the Brae equation that the half-height width becomes smaller and the grain size becomes larger. What’s more, the grain size can be estimated from the Scherrer equation, and it can be found that the grain size increased as the FWHM increased [25]. This is because the preheating of the substrate caused the fluctuation of the molten pool size, and affected the solidification conditions of 24CrNiMoY alloy steel powder at the solidification interface, making the primary dendrite arm spacing (PDAS) fluctuate within a certain range [16], which increased the grain size.

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The algorithm for basic calculation of grain thickness using FWHM is as follows:

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Firstly, the value of FWHM was converted into a radians value in this paper, and the formula is:

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FW（S） FWHM 

180

(2)

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Then, the Scherrer formula [26] was used to calculate the grain thickness, the

K FW ( S )  cos

(3)

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D

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formula is:

wherein D represents the mean grain size (nm), K is a constant, generally 0.89, λ is the

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wavelength of the X-ray (nm), and the X-ray wavelength of the copper target is 0.154 nm, FW(S) indicates that the sample is broadened (Rad). In addition, in the DLD process, the nucleation rate n depends on the driving force of crystal formation, and this driving force is the degree of subcooling ΔT. The relationship between n and ΔT is expressed by a Gaussian equation as shown in Eq. (4) [27]: T

n(T ) 

n dn 0 d (T )d (T )  max T

T

 exp[ 0

(T  TN ) 2 ]d (T ) 2(T ) 2

(4)

In the formula (8), ∆TN represents mean supercooled, ΔTσ is standard deviation, nmax indicates maximum nucleation density. Theoretically, analysis of Eq. (4) shows that ΔT decreased as the preheating temperature increased, which led to a decrease in n(ΔT) and grain coarsening in the process of DLD 24CrNiMoY alloy steel samples,.

Journal Pre-proof But in fact, the average grain size of the three samples (Non-preheating, 200℃ and 250℃) obtained by Channel 5 software was reduced from 1.2 μm without preheat treatment to 0.99 μm (200℃), and increased to 1.95 μm when preheated to 250℃, as shown in Fig. 7(a), (c) and (e). This phenomenon does not match Eq. (4), the reason for this phenomenon was that the proper preheating temperature (200℃) reduced ΔT, while the high temperature environment by preheating satisfied the formation temperature of GB, thereby starting to form a large amount of GB and improving n(ΔT), which causing grain refinement. When the preheating temperature reached

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250℃, the heat of the sample was accumulated during the layer-by-layer deposition

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solidification process so that the heat did not reach the diffusion, thereby causing the growth and sudden coarsening of the crystal grains.

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It can be seen from the grain boundaries of Fig. 7(b), (d), and (f) that as the

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preheating temperature changed, the content of the small-angle grain boundary (red

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line) was significantly reduced, which indicates significant coarsening of the structure, and the proportion of high-angle grain boundary (black line) increased from 29.0% to

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40.4% when the preheating temperature was at 200℃ compared with the non-preheating sample. The high-angle grain boundaries generally corresponded to

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lath-like columnar crystals and equiaxed grain boundaries were formed by some highly undercooled regions. Their presence helps to improve the properties of the samples, meanwhile they can inhibit or even prevent the propagation of cracks and improve the toughness of the sample.

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Fig. 7. EBSD maps and grain boundaries of the samples fabricated at different preheating temperature: (a,b)Non-preheating, (c,d)200℃, (e,f)250℃.

Fig. 8 illustrates the cross-section structure SEM images of the samples at different preheating temperature, the structure was mainly bainite. Among them, bainite was composed of granular bainite (GB) and lath bainite (LB), and bainite of two forms were interlaced. As can be seen in Fig. 8(b) and (c), at lower preheating temperatures (100℃ and 150℃), the structure was mainly laminated LB. With the preheating temperature increased, the cooling rate slowed down, the grains were gradually coarsened, most of the austenite were transformed into GB, and the content of LB was reduced by 24% from 100℃ to 200℃, as shown in Fig. 8(f). At this time (250℃), the main structure of the sample was GB (Fig. 8(e)). The GB had two

Journal Pre-proof structure forms [28], one was polygonal ferrite (PF) and carbon-rich austenite islands which was unevenly distributed around it, such a GB was called GB1, as shown in Fig. 9(a). The other was composed of short strip ferrite formed by the segmentation of LB and surrounding carbon-rich austenite, which was called GB2, as shown in Fig. 9(b). The island structures in GB1 and GB2 were derived from the carbon-rich residual austenite region, but their distribution and morphology in the ferrite matrix were completely different, which was closely related to the ferrite during the transformation

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[21].

Fig. 8. SEM micrographs showing the representative microstructure of the deposited 24CrNiMoY alloy steel with different preheating temperature: (a)Non-preheating, (b) 100℃, (c) 150℃, (d) 200℃, (e) 250℃, (f) Bainite content under different preheating parameters.

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Fig. 9. Morphology of different granular bainite: (a)GB1, (b)GB2.

Fig. 10(a) shows the TEM morphology of the sample when the substrate was not

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preheated, it can be seen that the ferrite slabs in the LB and the average width of the

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slats was 280nm. Fig. 10(b)-(d) show the TEM structure at 200℃ preheating, the high temperature due to the preheating was higher than the transformation temperature of

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the bainite, thereby causing the formation of PF and the precipitation of rod-shaped

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carbides (Fig. 10(b)). It can be seen in Fig. 10(c), there was a large amount of striped

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structure on the ferrite matrix. According to the diffraction spot analysis, the striped structure in Fig. 10(c) was ferrite nano twin. The nano twin structure was a typical

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plastic phase, which could increase the plasticity and increase the strength of the alloy [29]. The formation of a small amount of nano twins was caused by complex thermal

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stresses in the DLD process and lattice distortion during the forming process. Fig. 10(d) shows the GB in which the island structure was unevenly distributed. This was due to the growth of ferrite around the carbon-rich island structure. During the solidification transformation, with the change of heating temperature and holding time, the ferrite matrix was gradually formed in the carbon-depleted region, and the area of the carbon-rich austenite in the ferrite matrix was gradually reduced, forming the carbon-rich island-like structure with irregular distribution.

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Fig. 10. TEM image without preheating and preheating at 200℃: (a)LB, (b) Cementite in bainite, (c) Twinning, (d) GB.

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4. Mechanical properties

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4.1 Microhardness of DLD 24CrNiMoY alloy steel samples Fig. 11 shows the microhardness profile along the cross section of as-deposited

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samples (about 5mm in height). As can be seen in Fig. 11(b), the mean microhardness of the samples with different preheating temperature of 25℃, 100℃, 150℃, 200℃

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increased from 354 HV0.2 to 414 HV0.2, however when preheated to 250℃, its microhardness was significantly decreased to 344 HV0.2. This is mainly attributed to the structure and defect of the deposited layers. According to Fig. 7, the structure of the deposited layer was a mixture of GB and LB at 200℃, while it was mainly GB at 250℃. Besides, as the preheating temperature increased, the microstructure distribution in the deposition layer was more uniform, so the average hardness was also slightly improved [17].

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Fig. 11. Microhardness distribution of deposition layer under different preheating temperatures: (a)deposition layer from top to bottom, (b)hardness distribution.

4.2 Tensile behavior

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The stress-strain curves of the five sets of samples at different preheating

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temperatures are shown in Fig. 12. Fig. 13 shows the ultimate tensile strength (UTS), yield strength (YS) and elongation (EL) of the five samples. It can be deduced from

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Fig. 13 that the tensile behavior and hardness of the sample appear a similar change

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trend. When preheating at 200℃ the tensile behavior of the sample was optimally matched, its tensile strength and elongation were 1051 MPa and 7.4%, respectively.

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This was due to the mixed structure of the sample with GB (34.7%) and LB (65.3%) at this time. LB exhibits higher strength and poor toughness, GB expresses lower

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strength and good toughness, a certain amount of mixed structure of GB and LB can make the sample have both strength and toughness. At 250℃, its strength did not change much (the strength still maintains around 1000 MPa) while the elongation was significantly reduced by 3.3%. In theory, the strength of the sample at 250℃ should be lower than the other samples, its elongation after fracture was better than the other samples at different preheating temperatures, but the actual test results were different. This was due to the fact that the formability in samples, morphology of the bainite and grain size played an important role in the tensile property [30].

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Fig. 12. Stress-strain curve of 24CrNiMoY deposited layer.

Fig. 13. Tensile properties of 24CrNiMoY alloy steel samples at room temperature.

Fig. 14 is the SEM images of tensile fracture of non-preheated and preheated samples at 100-250℃, all of which were ductile fractures. As shown in Fig. 14(a) and (b), the fracture surfaces were mainly some small dimples and a small proportion of dissociation surface, and there were micropores at the bottom of the dimples, the toughness of the samples was relatively low. According to Fig. 14(c), the fracture of the sample appeared as uniformly distributed shallow dimples, and the toughness of the sample was slightly increased compared to the samples with non-preheating and preheating at 100℃. However, the fracture morphology of Fig. 14(e) was mainly parabolic dimples and shallow dimples, so the toughness was markedly reduced from

Journal Pre-proof 7.4% to 5.2%. As can be seen in Fig. 14(d), a large number of larger equiaxed dimples were observed in the sample at the preheating temperature of 200℃. Compared with the shallow dimples and parabolic dimples in other samples, the equiaxed dimples indicated a better toughness, thus the toughness of the samples was best when the preheating temperature was 200℃. This is mainly due to the morphology and grain size of the bainite in the deposited alloy steel sample at this time, the structure of the sample were 65.3% LB and 34.7% GB (Fig. 8(f)), the average grain size was refined

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with the value of 0.99 μm (Fig. 7(d)).

Fig. 14. Fracture profile of 24CrNiMoY tensile specimens with different preheating temperatures: (a)Non-preheating, (b)100℃, (c)150℃, (d)200℃, (e)250℃.

5. Discussion We studied the generation mechanism of bubbles in the molten pool during the

Journal Pre-proof DLD process and the evolution of deposited structure, as well as their impacts on the performance. These mechanisms are discussed below as they are related to the solidification process normally presented during DLD. 5.1 Effect of preheating temperature on crack and porosity The formability of the samples after appropriate preheating treatment was significantly improved due to the thermal stress during the DLD process. In laser direct deposition alloy steel, residual stress is one of the main factors leading to cracks. In fact, there are two factors that dominate the residual stress level of the deposited

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layer, one is the difference in thermal expansion coefficient between Q235 and

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24CrNiMoY, and the other is the temperature difference between the deposited layer and the substrate. In the DLD process, both factors can affect the residual stress level

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by changing the temperature field during solidification. Among them, the temperature

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difference between the substrate and the deposited layer is the dominant factor. This

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difference in temperature can lead to a difference in cooling rate, which in turn causes the shrinkage to be out of sync, thus determining the level of tensile stress in the

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solidification process. On the other hand, when the cooling is nearing the end, the temperature difference is already very small, yet the deposited alloy steel layer and

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the substrate still have a certain internal stress due to the unsynchronized shrinkage. These internal stresses will generate different stresses in different directions (X, Y, Z) of the deposited layer, resulting in different plastic strains. Cracking occurs when the plastic tensile deformation in a certain direction is greater than the maximum allowable deformation of the material [31]. Therefore, when the substrate is preheated at a certain temperature, the temperature gradient between the substrate and the deposited layer could be significantly slowed down, thereby reducing the residual stress of the deposited layer during solidification and preventing the generation of cracks as shown in Fig. 3. The presence of pores in samples not only reduces the formability, but also affects its mechanical properties. The formation and evolution of voids mainly depend on the formation and evolution of bubble in laser molten pool. The movement of the bubble in the molten pool depends mainly on the solidification state of the molten

Journal Pre-proof pool. When the rate of bubbles overflowing the molten pool in the molten pool is lower than the rate of solidification of the metal, pores are generated, and the Eq. (5) [32] is the condition for forming the pores: Ve  V0

(5)

Where Ve is the buoyancy speed of the bubble, V0 is the solidification speed of the molten pool. When Ve < V0, pores inside the deposited layer are formed; when Ve = V0, pores on the surface of the deposited layer are formed; when Ve > V0, no pores are

2(    e ) gl 2 9

(6)

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Ve 

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generated. The bubble floating speed is represented by the Stocks formula [33]:

-p

Where η is the viscosity of the liquid metal (Pa·S), ρ is the liquid metal density (kg/m3), ρe is the bubble density (kg/m3), g is the gravitational acceleration (m/s2), and

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l is the bubble radius (m). It can be seen from the Eq. (6) that the bubble is subjected

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to η. When η is larger, the bubble floating speed is slower, and the bubble is less likely to escape the molten pool, thereby forming pores remaining in the sample. However,

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η is mainly affected by the temperature, η decreases as the temperature increases, and η increases remarkably when the liquid metal solidifies and crystallizes.

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Pores are formed when bubbles do not escape from laser molten pool, so the formation of pores is attributed to the generation of bubbles. However, the causes of bubbles are different under different processes. Kang et al. [34] studied the cause of pore formation by laser cladding alloy powder with oxygen content of 0.53%, and formed Cr2O3 which is more stable than CO and CO2 by adding Cr particles to reduce bubble formation and prevent pore formation, so a non-porous alloy steel was obtained. Cao et al. [22] established the pore behavior model by studying the changes of pores at different energy densities, and explained the causes of pore production. Therefore, one of the causes of pore formation in the 24CrNiMoY alloy steel sample is that the oxides in the high temperature molten pool reacted with C to form CO and CO2 bubbles. Besides, the flow of shielding gas in the DLD process had an effect on the pores. Chen et al. [35] studied the effect of different process parameters on the

Journal Pre-proof formation of pores. The results showed that the shielding gas would be entrapped with the beam, forming bubbles in the depression of the molten pool, and forming pores when it cannot escape from the molten pool. According to Fig. 2, the shielding gas can be sent to the molten pool under the action of the powder feeding head to form bubbles. The shielding gas is entangled in the deposited layer along with the powder and the laser beam. The bubbles are formed in the depressed portion of the molten pool, and cannot remove from the molten pool to form pores. In addition, 3% of the hollow spheres [36] present in the alloy steel

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powder cause the gas (Ar) contained therein to enter the molten pool during the DLD

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process, and the gas aggregates to form bubbles under the action of the molten pool. When the bubbles in these molten pools cannot escape in time before solidification,

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the formed pores remain in the deposited sample.

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Based on previous studies [22, 34-36], the distribution of the pores mainly

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depended on the convection motion of the melt, and the convection motion of the melt was determined by the temperature of the molten pool. The uneven distribution of the

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surface temperature of the molten pool would create a tension gradient on the surface of the molten pool, which caused the melt to flow from a portion with a low surface

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tension to a high portion [37]. The relationship between surface tension gradient and temperature satisfied the following Eq. (7) [38]: d T dT  c(1  ln ) dx T0 dx

(7)

Where dσ/dx represents the surface tension temperature coefficient of the molten pool, dT/dx stands for the molten pool temperature gradient, c is the specific heat capacity. At the same time, the surface tension gradient has the following relationship with the convection motion acceleration of the molten pool [39], that is: d   dx

(8)

Where ρ means the density of the metal, Δ indicates the thickness of the liquid layer participating in the flow, α denotes the acceleration of the convection of the molten pool. Substituting the Eq. (7) into the Eq. (8) and obtaining the Eq. (9), the

Journal Pre-proof acceleration of the convection of the molten pool is: 

1





d c T dT  (1  ln ) dx  T0 dx

(9)

The formability of the five samples (with and without preheating) had a different morphology probably because of the different molten pool temperature and solidification rates caused by the different preheating temperatures. Fig. 15 shows a schematic diagram of the formation mechanism of the pores in the DLD process, it can be seen that the pores in the sample mainly underwent three processes of

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formation, floating, and growth. As shown in Fig. 4(a), the pores remained in the

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sample without preheating, this is because the bubble formed by the shielding gas which was entangled in the molten pool and the bubble which was accumulated by the

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metallurgical reaction did not escape, meanwhile the solidification rate of the molten

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pool was fast, the liquid pool maintained a short liquid time, therefore the bubble was

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not easily escaped from the molten pool and remained in the sample to form pores. But when starting preheating and the preheating temperature was gradually increased, the temperature of the molten pool increased, melting pool solidification

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rate slowed down. According to Eq. (6) and (9), η was decreased and the convection

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acceleration of the molten pool was increased, which was favorable for the floating and overflow of the bubbles (Fig. 14(b) and (c)). Therefore, when the preheating temperature was 200℃, there were almost no pores in the sample or pores with extremely small size remained, an equilibrium state in which the molten pool was well solidified and free of bubbles was achieved. However, when the preheating temperature was up to 250℃, as the surface temperature of the molten pool increased further, the vapor pressure on the metal surface became larger, thus the metallurgical reaction inside the molten pool was more intense, which led to the more protective gas trapped in the molten pool. At the same time, when the diffusion and solidification process was prolonged, and small bubbles combined to form large bubbles, thus more pores were formed after solidification was completed, as shown in Fig. 4(e) and Fig. 15(d).

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Fig. 15. Schematic diagram of pores formation mechanism: (a)Start forming pores, (b)Pores floating, (c)Pores spilling, (d)Pores grow up.

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5.2 Effect of preheating temperature on microstructure evolution

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Fig. 16 shows the evolution mechanism between the bainite structure and preheating temperature of 24CrNiMoY alloy steel. It can be seen in Fig. 16(a) that the

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structure of the sample was mainly LB when the substrate was not preheated, and the formation temperature of LB was low. This is because that first ferrite nucleus was

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formed on the austenite grain boundary or the carbon-poor region that near the grain boundary, which grew parallel from the austenite grain boundary into the crystal, forming a parallel ferrite lath bundle (Fig. 16(a)). At the same time, there was a carbon-rich austenite arranged in a direction substantially parallel to the lath ferrite, and these carbon-rich austenite particles were fine and ordered. As the substrate began to be preheated and the preheating temperature was slowly increased, the temperature of the molten pool during the DLD deposition process was also gradually increased, and the formation temperature of the GB was higher than that of the LB, thus GB was formed, as shown in Fig. 16(b). During the formation of GB, the ferrite precipitated from the supercooled austenite was initially lath-like and had a high dislocation density. However, since the DLD was deposited layer by layer and the center temperature of the deposited layer

Journal Pre-proof increased with the preheating temperature, the ferrite would recover at such a higher temperature, the slats were not as obvious as the LB and gradually appeared PF. The GB was composed of PF and randomly distributed carbon-rich austenite islands on the ferrite matrix. It can be seen in Fig. 16(b) that these small amounts of disorderly distributed GB had a segmentation effect on the LB, so that the long pieces of LB was divided into short strip LB. As the preheating temperature continued to rise and the cooling rate slow down, a large amount of GB was formed, the LB was gradually reduced due to the division, and the deposited layer structure was mainly GB, as

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shown in Fig. 16(c) and Fig. 8(e).

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Fig. 16. Evolution mechanism between bainite morphology and preheating temperature of 24CrNiMoY alloy steel: (a)Non-preheating, (b)200℃, (c)250℃.

5.3 Effect of preheating temperature on properties In addition to the microstructure, porosity defects in the sample can significantly affect hardness and tensile behavior. According to Wakshum et al. [40], the porosity defects in the sample had a great influence on the microhardness of the sample. Therefore, the analysis of Fig. 4 and Fig. 5 could explain why the average microhardness was the highest at 200℃. Another reason for explaining the hardness was based on the grain size analysis of Fig. 7. Jahangiri et al. [41] studied fine grains and pointed out that the reduction of porosity content means better microhardness. However, when the preheating temperature was too high, the microhardness value was significantly lowered. Therefore, by applying a suitable preheating temperature to

Journal Pre-proof the substrate, a sample having a high hardness and a uniform distribution of hardness values can be obtained. The explanation about the tensile property is mainly divided into the following two aspects. The first explanation was due to the dispersion strengthening of carbon-rich austenite islands. Since the islands were hard phases, and they were precipitated in a small dispersion, interacting with dislocations, thereby hindering the movement of dislocations, finally increasing the strength of the steel by means of diffusion strengthening. Meanwhile an increase in the strength of the hardened phase

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or an increase in the volume fraction caused a decrease in the yield ratio, thereby

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improving the tensile strength without affecting the yield strength [42]. In addition, when the preheating temperature was 250℃, the deposited layer structure was GB

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mixed with GB1 and GB2. GB2 had the properties of LB and also had high strength,

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so it has high tensile strength and low yield strength at this time. Secondly, it was

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explained from the aspect of elongation because the grain size also had an important influence on the strength of the sample. Theory analysis and the experimental results

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indicate that the smaller ferrite grain size is, the higher the yield-strength ratio is [43]. According to the Hell-pitch formula, the smaller the grain size of the metal member,

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the better the strength, and the better the toughness. The sample had the largest average grain size (1.99 μm) at 250℃ (Fig. 7), thus its mechanical properties were worst. Moreover, the defects of the sample would also have an effect on the tensile properties, as shown in Fig. 3(e) and Fig. 4(e). Strong plastic product is one of the key indicators to judge whether the material strength and toughness match. Therefore, according to Fig. 13, the sample with preheating temperature of 200℃ has the highest strength and plasticity product, indicating that the mechanical properties of strong toughness matching can be obtained under this parameter. Thus, although preheating could uniform the structure and reduce thermal stress, the unsuitable preheating temperature would reduce the mechanical properties. The preheating temperature had a significant effect on the final structure and mechanical properties of DLD 24CrNiMoY alloy steel samples. Three different types of bainite could be obtained at different preheating temperatures,

Journal Pre-proof among which LB could improve the strength of the alloy steel, and the toughness and plasticity of GB were better than that of LB. Therefore, by controlling the preheating temperature, the proportion of LB and GB in the steel could be adjusted, thereby obtaining a sample with better comprehensive mechanical properties.

6. Conclusions (1) Proper preheating temperature can eliminate cracks and control porosity defects. During the deposition of 24CrNiMoY alloy steel samples, the temperature

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gradient between the deposited layer and substrate can be significantly reduced

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by preheating substrate, thereby effectively reducing the generation of cracks. In addition, as the preheating temperature increased, the amount of pores in the

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sample decreased first and then increased, the sample structure was the densest

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when the preheating temperature was 200℃, and there were almost no pores.

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(2) Preheating has a certain regulation effect on the microstructure of laser deposited 24CrNiMoY alloy steel samples. The main structure in the five groups of samples

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was a mixed structure of LB and GB. As the preheating temperature increased to 200℃, the proportion of LB decreased from 95.7% to 65.3%, while the content of

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GB changed from 4.3% to 34.7%. As the preheating temperature increased to 250℃, the amount of GB reached a maximum of 96.5%. (3) The suitable preheating temperature has the effect of refining the structure of the deposited alloy steel samples. The grain size decreased from 1.20 μm to 0.99 μm at first and then increased to 1.95μm with the rise of preheating temperature. Besides, as the preheating temperature increased to 200℃, the proportion of high-angle grain boundaries increased from 29.0% to 40.4% and that of low-angle grain boundaries decreased from 71.0% to 59.6%, the change of high-angle grain boundaries revealed the promotion at property under the preheating temperature of 200℃. (4) The suitable preheating temperature has an important influence on the hardness, tensile strength and toughness properties of DLD 24CrNiMoY alloy steel samples. With the rise of preheating temperature (25℃, 100℃, 150℃, 200℃ and 250℃), the

Journal Pre-proof average hardness value of the deposited layer first experienced an increase from 345 HV0.2 to 414 HV0.2, and then decreased by 70 HV0.2 at 250℃. At 200℃, the microhardness of the deposited layer was the most uniform and highest, meanwhile the tensile behavior of the sample was the best match of strength and toughness, its UTS and YS were 1051 MPa and 915 MPa, the elongation was 7.4%. Acknowledgments

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This work was financially supported by National Key R&D Program of China (2016YFB1100201), Green Manufacturing System Integration Project of the Industry

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and Information Ministry of China (2017), Research and development plan for the

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future emerging industries in Shenyang (18-004-2-26), and Shenyang Achievement

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Transformation Project (2019). Data availability statement

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The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

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References

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Graphical Abstract (b)

(a)

(c)

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Fig. 1. Schematic diagram of DLD process ma and the evolution mechanism between preheating

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and pores and structure:(a)The DLD process diagram, (b)Schematic diagram of pores formation

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mechanism, (c)Evolution mechanism between bainite morphology and preheating temperature of

Abstract:

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24CrNiMoY alloy steel

The direct laser deposition (DLD) process, evolution mechanism between

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preheating temperature and structure are visually described in this graph. Fig. 1(a) displays the DLD process design for preheating the substrate, the main heat transfer

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method is heat conduction. Fig. 1(b) shows the formation mechanism of the pores in the DLD process, it can be seen that the pores in the sample mainly underwent three

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processes of formation, floating, and growth. Fig. 1(c) exhibits the evolution mechanism between the bainite structure and preheating temperature of 24CrNiMoY alloy steel, it can be deduced the microstructure of the sample changes from lath bainite (LB) to granular bainite (GB) with the rise of preheating temperature. The proper preheating temperature of 200℃ can effectively eliminate the crack, decrease the pores and significantly improve the structure and properties of the DLD 24CrNiMoY alloy steel.

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Highlights  Preheating significantly eliminates cracks.  The effect of preheating temperature on pore formation was analyzed.  The mechanism for preheating temperature and bainite evolution was established.

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 The optimal preheating temperature of low carbon alloy steel in DLD was

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na

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obtained and optimal structure and performance were prepared.

Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Figure 6

Figure 7

Figure 8

Figure 9

Figure 10

Figure 11

Figure 12

Figure 13

Figure 14

Figure 15

Figure 16