Evolution in microstructure and high-temperature oxidation behaviors of the laser-cladding coatings with the Si addition contents

Journal Pre-proof Evolution in microstructure and high-temperature oxidation behaviors of the lasercladding coatings with the Si addition contents Y.L. Zhang, J. Li, Y.Y. Zhang, D.N. Kang PII:

S0925-8388(20)30494-1

DOI:

https://doi.org/10.1016/j.jallcom.2020.154131

Reference:

JALCOM 154131

To appear in:

Journal of Alloys and Compounds

Received Date: 9 December 2019 Revised Date:

21 January 2020

Accepted Date: 31 January 2020

Please cite this article as: Y.L. Zhang, J. Li, Y.Y. Zhang, D.N. Kang, Evolution in microstructure and high-temperature oxidation behaviors of the laser-cladding coatings with the Si addition contents, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154131. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2020 Published by Elsevier B.V.

Y.L. Zhang: Experiment, Analysis, Writing- Original draft preparation. J. Li: Supervision, Methodology. Y.Y. Zhang: Software, Validation, Writing- Reviewing and Editing. D.N. Kang: Software, Validation.

Evolution in microstructure and high-temperature oxidation behaviors of the laser-cladding coatings with the Si addition contents Y.L. Zhang, J. Li*, Y.Y. Zhang, D.N Kang School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China

Abstract In view of poor oxidation resistance of Ti6Al4V, the composites coatings prepared by laser cladding were used to solve the above shortcoming with a mixture of NiCrBSi and 5 wt.%- 30 wt.% Si as the cladding material. The investigation into the evolution in microstructure, mechanical properties and oxidation resistance with the Si content was carried out in detail. The results indicated that the coating with a Si content of 5 wt.% was mainly composed of Ti2Ni, TiNi, TiB2, TiB and TiC. A new phase of Ti5Si3 was synthesized in the coatings with a Si content more than 10 wt.%, and its content presented an upward tendency with the Si content coupled with the decrease in TiB2, TiB and TiNi. The microstructural change caused the slight increase (about 3.82%) in microhardness of the coatings, accompanied with which their fracture toughness was correspondingly reduced by about 23.90%. The increase in Si content significantly improved oxidation resistance of the coatings due to the reduction in oxidation weight gain. The oxidation weight gain of the coatings was reduced about 36.98% when the Si content was increased from 5 wt.% to 30 wt.%. A compacter oxidation layer could be formed on the coating’s surface with a higher Si content due to the higher PBR （Pilling-Bedworth ratio）value for Ti5Si3 subject to oxidation than those of the other phases, which played the essential role in the improvement in the coatings’ oxidation resistance. The suitable addition of Si was

confirmed as 20 wt.%, in which the coating was endowed with excellent comprehensive mechanical properties and outstanding oxidation resistance.

Keywords: Titanium alloy; laser cladding; coatings; oxidation resistance; fracture toughness

1. Introduction Titanium alloys are increasingly applied in various fields, such as biomedical, aerospace, and automotive industries. This should be attributed to their high specific strength (strength-weight ratio), outstanding corrosion resistance, and excellent biocompatibility [1-3]. However, there are few reports on the use of titanium alloys at the temperatures high than 600 °C, since they are subject to serious oxidation in such a high-temperature environment. This will greatly shorten its service life. Surface modification is considered as an effective method to solve the above shortcoming involved in titanium alloys, by which a coating with outstanding oxidation resistance is fabricated on their surface. At present, many techniques have been explored and applied in surface modification of titanium alloys, which include laser remelting [4], laser cladding, nitriding [5], carburizing [6], plasma spraying [7], chemical vapor deposition [8], physical vapor deposition [9], etc. Comparatively speaking, laser cladding as a new and promising surface modification method demonstrates several advantages, such as (1) obtaining a dense structure, (2) not being limited by the type of substrate material and cladding layer; (3) having high energy density, no pollution, and being easy to realize [10, 11]. Some laser-clad had been prepared to enhance poor resistance to oxidation of titanium alloys. Maliutina et al. [12] prepared a TiAl-based coating on Ti6Al2Sn4Zr2Mo by laser cladding Ti48Al2Cr2Nb powder. The coating exhibited a better oxidation resistance than the substrate at 700-900 °C for 100 h. This should be attributed to the

formation of dense Al-rich oxides on the coating’s surface. Moreover, the presence of Nb and Cr also played a positive role in inhibiting the growth of TiO2. Jenny Cecilia Zambrano Carrullo et al. [13] also applied Ti48Al2Cr2Nb powder to produce a protective coating on Ti6Al4V by laser cladding, and further confirmed that the Ti48Al2Cr2Nb coating had better oxidation resistance than Ti6Al4V in air at 800 °C for 5, 10, 25, 50, 100 and 150 h. Feng et al. [14] fabricated a (Ti3Al+TiB)/Ti composite on Ti6Al4V by laser cladding to improve its high temperature oxidation resistance. The coating showed a better resistance to high temperature oxidation than the substrate due to its lower weight gains (only about 20-30 those of the substrate). The existence of Al2O3 was responsible for the change, which made the oxidation layer denser and inhibited the diffusion of oxygen. Liu et al. [15] produced a TiN/Ti3Al laser-clad coating on Ti6Al4V, and evaluate its oxidation resistance at 600 °C and 800 °C, respectively. The oxidation resistance of the coating was highly higher that of the substrate. The relative oxidation resistance of the coating at 600 °C and 800 °C reached 6.83 and 1.94, which was far higher than that of the substrate (1). The outstanding resistance to high temperature oxidation mainly resulted from the formation of TiN, Al2O3 and TiO2. With respect to this technology, the cladding material is regarded as an essential factor deciding the laser-clad coatings’ final properties. The self-fluxing alloy as a cladding material is widely employed. According to the main components, the selfmelting alloy can be divided into nickel-based, iron-based, cobalt-based, and other self-fluxing alloys. Nickel-based alloy powders are most commonly employed because of their advantages of good wettability, good corrosion resistance, and hightemperature self-lubricating action and a good metallurgical bond with the titanium alloy matrix during laser cladding [16, 17]. Many investigations into the laser-clad

coatings’ oxidation behaviors with nickel-based alloy as main component have been carried out. Guo et al. [18] fabricated NiCrBSi and NiCrBSi/WC-Ni composite coatings on pure Ti by laser cladding. Their high temperature oxidation behaviors were investigated from 25 to 1100 °C with pure Ti as a reference material. The results indicated that the two coatings demonstrated a better oxidation resistance than pure titanium. Ti was oxidized at a temperature of about 300 °C, which was significantly lower than that of the laser-clad coatings (about 650 °C). Moreover, the oxidation of Ti entered into a fast stage at a lower temperature of about 500 °C when compared with the coatings (about 750 °C). Liu et al. [19] produced a laser-clad coating on TA2 titanium alloy with the Ni based alloy powder (GH4169 powder) as the cladding material, and investigated its oxidation behaviors. NiTi, NiTi2 and Ni3Ti were synthesized as the main phases in the laser-clad coating. The coating exposed at 800 °C for 100 h demonstrated more excellent oxidation resistance than TA2 due to better oxidation resistance of NiTi, NiTi2 and Ni3Ti than pure titanium. Moreover, a dense layer of Al2O3 resulting from the oxidation of Al involved in the coating also contributed to improving the coating’s oxidation resistance. Aghili et al. [20] fabricated a composite coating on titanium aluminide by cladding the NiCr-Cr3C2 powder mixture. The coating was mainly composed of γ solid solution reinforced by Cr3C2, Cr7C3 and Cr23C6. The oxidation testes indicated that the specific weight gain of the coating was about 0.004 mg•cm-2, which was twice lower than that of the substrate (about 0.009 mg•cm-2). This should result from the formation of Cr2O3 as the main phase on the coating’s surface, which acted as a barrier to slow the diffusion of oxygen into the coating. Lv et al. [21] investigated the effect of the TaC addition content on oxidation resistance of the TiNi/Ti2Ni matrix coatings. The addition of TaC significantly improved the coatings’ oxidation resistance exposed at 800 °C for

50 h due to the oxidation weight gain reduced from 3.30 mg•cm-2 to 1.69 mg•cm-2 with the addition of TaC increased from 0 wt.% to 40 wt.%. This change resulted from the formation of Ta2O5 in the oxide layer adhering to the surfaces of these coatings with TaC. Xu et al. [22] produced a protective laser-clad coating on Ti6Al4V with the Ni80Cr20-40Al-20Si (mass fraction, %) precursor mixed powders as the cladding material. The coating mainly consisted of Al3Ti/NiTi intermetallic compounds as the matrix, and Ti5Si3/Al3Ni2 as the reinforced particles. Ni and Al involved in the coating had lower affinity to oxygen than Ti in the substrate at 800 °C for 32 h, resulting in oxidation resistance of the coating superior to that of the substrate. It can be noticed that some metals and their compound are usually added to the nickel-based alloy powder to improve the final coatings’ oxidation resistance. The addition of Si element had been confirmed as an important factor enhancing oxidation resistance of different alloys prepared by the other methods. Tunthawiroon et al. [23] investigated oxidation behaviors of Co29Cr6Mo alloys with different Si contents exposed to air at the high temperatures of 700-1000 °C for 2, 4, 12, 24, 48, 72, and 100 h. The results demonstrated that the Si addition was beneficial to the formation of a continuous Cr2O3 film. Moreover, the formation of SiO2 inhibited the oxygen’s inward diffusion and chromium’s outward diffusion. Esleben et al. [24] compared oxidation resistance of two alloys with respect to Co17Re18Cr15Ni and Co17Re18Cr15Ni2Si alloy in the temperature range 800 °C-1100 °C. The investigations proved that the addition of Si played the positive role in the improvement of oxidation resistance due to the two reasons: (1) a discontinuous SiO2 layer retarded the volatilization of Re-oxides; (2) the formation of Si-oxide particles promoted the nucleation of Cr2O3 as the diffusion barriers slowing the transportation of metal atoms during oxidation. Guo et al. [25] also investigated the effect of the Si

addition on oxidation resistance of Ti0·5Al0·5N coatings. The alloying of Si resulted in a compact sublayer rich in Al2O3 easily formed, which improved the coatings’ oxidation resistance. The other investigations also confirmed the positive effect resulting from the addition of Si [26-30]. Unfortunately, present investigations into the effect of Si on the laser-clad coatings’ oxidation resistance were hardly reported. In this study, the laser-clad coatings were fabricated on Ti6Al4V. The NiCrBSi with different Si addition contents was employed as the cladding materials. The effects of the Si content on microstructure, mechanical properties and oxidation behaviors of the coatings were revealed in detail. The evolution in oxidation mechanism with the increase in Si content was emphasized. The suitable Si content was confirmed, in which the coatings demonstrated the excellent comprehensive mechanical properties and outstanding oxidation resistance.

2. Experimental procedures 2.1 Preparation of cladding layers and laser cladding In this research, Ti6Al4V with chemical compositions of 88.8 wt.%, 6.5 wt.% Al and 4.3 wt.% V was selected as the substrate. A Ti6Al4V round bar with a diameter of 50 mm was cut into the discs (Φ50 mm*10 mm), then the cladding surfaces of those discs were subject to milling and ultrasonic cleaning in acetone, so that the contaminates and oxides adhering to the surfaces could be completely removed. The nickel-based alloy powder (75Ni16Cr3.5B4.5Si in wt.%) and high-purity Si powder (99.999 wt.%) were used as the cladding materials. Six kinds of mixed powders with different ratios were prepared for laser cladding (5 wt.%Si+95 wt.% NiCrBSi, 10 wt.% Si+90 wt.% NiCrBSi, 15 wt.%Si+85 wt.% NiCrBSi, 20 wt.%Si+80 wt.%NiCrBSi. 25 wt.% Si+75 wt.% NiCrBSi, 30 wt.% Si+ 70 wt.% NiCrBSi). The

obtained coatings were designated as coating I, coating II, coating III, coating IV, coating V, and coating VI. A pre-prepared layer was first prepared by a modified bonding method prior to laser cladding. Compared with the traditional bonding method, the usage amount of the binder is greatly reduced, the thickness of the pre-prepared layer can be precisely controlled and its compactness is correspondingly enhanced. Those advantages can significantly improve the microstructural reproducibility of the coating, and enhance its final quality. In this method, a 4 vol.% polyvinyl alcohol binder was brushed on the substrate surface, then the sample was placed inside a ring mold (inner diameter 50.2 mm, height 10.8 mm). The dried and uniformly mixed powder is evenly replaced on the substrate surface covered with the binder to fill the remaining upper space of the ring). The powder was pressurized for 3 min at a pressure of 30 MPa to former a prepared layer with a thickness of 0.8 mm. Finally, the sample was dried at room temperature for 24 hours before laser cladding. An YLS-5000 fiber laser was applied to obtain the laser-clad coating. The processing parameters (output power: 3 kW, scanning speed: 5 mm•s-1, spot diameter: 6 mm) were optimized to obtain a highqualified coating with a good interfacial bonding, a uniform microstructure free of porosity and cracks. In addition, an appropriate amount of the substrates were melted and participated in the chemical reactions occurring in the molten pool to form the desired phases. 2.2 Microstructural characterization and mechanical properties testing A PANalytical X' Pert Pro X-ray diffractometer was employed to detect the coatings’ phases. In order to remove the adverse effect of the comparatively coarse coatings’ surfaces on final results, the coatings’ surfaces were ground and polished with the abrasive papers prior to the detection. A HITACHI S-3400 scanning electron

microscope (SEM) coupled with a GENESIS EDAX energy dispersive spectrometer (EDS) was used to observe the coatings’ microstructure. The observed coatings were ground and polished in turn with #200, #800 and #1100 abrasive papers. Then they were etched for 30 s in a corrosive agent composed of 5 ml H2O, 8 ml HNO3 and 10 drops of HF. A HXD-1000 TMSC/LCD Vickers microhardness tester was employed to analyze the microhardness distribution throughout the cross sections of the coatings. The 200 gf load was employed for 15 seconds. The microhardness was firstly measured at the zones with a distance of 0.1 mm from the coatings’ surfaces. Then subsequent tests were performed at the zones with the same distance from the former. In order to reduce the testing errors, an average value was obtained from three values measured at three separated zones with a same distance from the coatings’ surfaces. The fracture toughness (KIC) was evaluated by the Vickers indentation method, which was carried out on a HV-50 Vickers hardness tester with an applied load of 50 N for 10 s. A VHX-600K optical microscope (OM) was used to observe the morphology in terms of the indentation and generated cracks. 2.3 High temperature oxidation experiments In order to precisely evaluate high-temperature oxidation resistance of the coatings resulting from the Si content, it is necessary to take out a part of the coatings from the cladding samples to remove errors caused by the oxidation of the substrate during the oxidation process. The six surfaces of each intercepted part were sequentially polished to a mirrored state with 320#, 800#, 1500#, and 2000# abrasive papers. The oxidation tests were carried out in a SG2-3-10 resistance furnace with a testing temperature of 800 °C for 200 h. During the tests, the samples were taken out

every 1 h for weighing on a Sartorius BSA124S balance with a 0.01 mg accuracy. The oxidation weight gain per hour per unit area was further calculated. The elemental compositions and chemical valence states of the samples suffering from oxidation were determined with an ESCALAB 250XiX-ray photoelectron spectrometer. The cross-sectional morphology of the coatings subject to the oxidation was observed by SEM.

3. Results and Discussion 3.1 Microstructural characterization The XRD patterns of the coatings with different contents of Si (5, 10, 15, 20, 25, 30 wt.%) are shown in Fig.1. The d values of the X-ray diffraction peaks of Coating I can match well with those of the JCPDS standard cards related to TiNi (No. 03-0654572), Ti2Ni (No. 01-072-0442), TiB2 (No. 01-075-1045), TiB (No. 01-089-3922) and TiC (No. 3-065-8805). It is clear that TiNi, Ti2Ni, TiB2, TiB and TiC as the main phase are synthesized in Coating I. When the addition content of Si reaches and exceeds 10 wt.%, some new diffraction peaks are clearly observed at 34.94°, 36.79°, 37.74°, 41.25°, 42.83°, 66.84°, 73.97°, and 77.32°. The indexed result indicates that above peaks originate from the formation of a new phase (Ti5Si3, No. 01-089-3721). Moreover, the content of Ti5Si3 presents an upward tendency with increasing the Si content, accompanied with reduction in content of TiNi, TiB2 and TiB. When the Si content is higher than 20 wt. %, only Ti5Si3, Ti2Ni and TiC are involved in the coatings. It is worth to note that all in-synthesized compounds in the coatings contain Ti, which should be attributed to the melting of the substrate during cladding. During laser cladding, besides a portion of the reflected energy, the other is absorbed by the substrate (Ti6Al4V) and the cladding material (NiCrBSi), causing a thin layer of

substrate surface melted. Ti from the substrate will react with B, C, Ni and Si from the cladding materials, resulting in the formation of Ti-rich compounds. The formation possibility of different compounds rich in Ti in this system can be predicted by thermodynamic calculations. Ti and Ni react to form compounds such as TiNi and Ti2Ni. Ti and B react to form compounds such as TiB and TiB2. Ti and C react to form compounds such as TiC. Ti and Si react to form compounds such as TiSi and TiSi2, and Ti5Si3. These compounds may be in synthesized as follows: Ti+Ni=TiNi

(1)

2Ti+Ni=Ti2Ni

(2)

Ti+B=TiB

(3)

Ti+2B= TiB2

(4)

Ti+C=TiC

(5)

Ti+Si=TiSi

(6)

Ti+2Si=TiSi2

(7)

5Ti+3Si=Ti5Si3

(8)

Based on thermodynamic data provided in the references [31], the relationship between change in standard Gibbs free energy (∆G0) of these reactions and temperature can be established. As illustrated in Fig. 2, the ∆G0 values in terms of all reactions are less than zero at the temperature ranging from 298 K to 1200 K, suggesting that the corresponding products can be spontaneously synthesized from the thermodynamic point of view. In terms of the compounds composed of Ti and Si, the preferential formation order is as follows: Ti5Si3﹥TiSi2﹥TiSi. This explains why Ti5Si3 is formed in the coating instead of the other two phases. When the high-power

laser is applied to the cladding materials, the cladding materials coupled with a layer of substrate will be melted. The molten pool composed of Ti, Ni, Si, B and C will be formed. During the subsequent cooling, TiB2 and TiC with higher molten points (3225 °C for TiB2, 3073 °C for TiC) will nucleate and grow firstly, followed by TiB (2200 °C) and Ti5Si3 (2130 °C). Finally, Ti2Ni (984 °C) and TiNi (1310 °C) with the comparatively low molten points are synthesized. Chemical compositions are not uniformly distributed within the whole pool due to Marangoni flow effects. With respect to Ti and B, some zones rich in B may be transformed into TiB2. Along with the gradual consumption of B resulting from the formation of TiB2, TiB with less content of B may be precipitated subsequently in some zones. According to the Ti-B binary diagram, another compound (Ti3B4) composed of Ti and B may be synthesized by a peritectic reaction between the liquid and TiB2 in a very narrow compositiontemperature range. Ti3B4 is not observed in the XRD patterns, which may be attributed to the rapid melting and rapid cooling characteristic of laser cladding. Some researches had also confirmed this [32, 33]. However, only the compound of TiC can be formed in the Ti-C system according to the Ti-C binary phase diagram. As far as the Ti-Si system is concerned, five compounds composed of Ti and Si can be formed, corresponding to Ti3Si, Ti5Si3, Ti5Si4, TiSi and TiSi2. Ti3Si, Ti5Si4 and TiSi are synthesized by the peritectic reactions, so it is very difficult to obtain them by laser cladding with a rapid solidification rate. TiSi2 with a high atomic ratio of Si and Ti is also difficult to be formed since the content of Si is less than Ti in the molten pool. Therefore, only the Ti-Si compound of Ti5Si3 is obtained in the coatings. With the decrease in Ti content resulting from the formation of TB2, TiB, TiC, Ti5Si3, its content is in the range from 50 at.% to 67 at.% in the Ti-Ni system. TiNi and Ti2Ni are finally deposited according to the Ti-Ni binary diagram.

Fig.3 indicates the cross sectional view of the whole laser-clad coating and the micrograph of the interface between the substrate and the coating. It is clear that the coating with a maximum thickness of about 2 mm is very dense (Fig.3 (a)). No obvious pores and cracks are observed. A continuous interface can be clearly observed the substrate and the coating, which indicates that a strong metallurgical bond is formed between the two (Fig.3 (b)). Clear inspection reveals that a transition zone with a thickness of about 350 µm is formed between the coating and the substrate. Fig.4 shows the microstructure from the cross sections of different coatings. For Coating I, some dark gray equiaxed dendrites, black blocky, and gray needleshaped particles are dispersed uniformly in the continuous gray phase (Fig.4 (a1)). The three kinds of scattered particles can be regarded as the reinforcements and the continuous gray phase can be taken as the matrix. The highly-magnified SEM image clearly reveals that the continuous gray phase contains two phases (a light gray convex phase and an off-white concave phase) (Fig.4 (a2)). The chemical compositions of the above-mentioned five phases (as marked in Fig.4 (a2)) were analyzed by EDS (seeing Table 1). In the continuous gray phase, the two phases are determined to be the secondary solid solution with the titanium-nickel compound as the solvent. For the light gray convex phase (zone 1), the contents of Ti and Ni are 41.52 at.% and 24.6 at.%, and the rest is a small amount of Cr, Al, V and Si. Since Cr, Al and V have the atomic radii similar to Ti (Cr: 1.85 Å, Al: 1.82 Å, V: 1.92 Å, Ti: 2.00 Å), these elements can replace the lattice position originally belonging to Ti in the solid solution. Similarly, Ni is replaced with Si due to their similar atomic radii (Ni: 1.62Å, Si: 1.46Å). The total atomic percentage of Ti, Cr, Al and V is 64.87 at.%, which is about twice that of Si and Ni (35.13 at.%). Combining with the XRD results, the phase can be confirmed as the Ti2Ni solid solution. With respect to the light gray

concave phase (zone 2), it can be determined to be the TiNi solid solution since the total atomic percentage (49.02 at.%) of Ti, Cr, Al and V is close to that of Ni and Si (50.98 at. %). The dark gray dendrites (zone 3) mainly contain Ti and C, which can be identified as TiC. The black blocky phase (zone 4) and the gray needle-shaped phase (zone 5) mainly consist of Ti and B, and can be considered as the boride. TiB with a structure of 27 is inclined to grow into the needle shaped particles due to its unit cell presenting a triangular prism. On the other hand, TiB2 has a crystal structure of C32, and its unit cell is a close-packed hexagon. Thus TiB2 demonstrates a massive morphology. Therefore, the black blocky phase (zone 4) and gray needle-shaped phase (zone 5) are determined as TiB2 and TiB, respectively. The area fraction of above-mentioned phases can be estimated (Fig.5 (a)). The total area fraction of the matrix is about 78.93%, in which TiN and Ti2Ni account for 23.68% and 55.25%, respectively. For the reinforcements (about 21.07%), the area fractions of TiC, TiB and TiB2 are 3.38%, 6.19% and 11.50%. When the Si content exceeds 10 wt.%, a new light gray phase is synthesized. EDS results indicate that the phase (zone 6 in Fig.4 (b2)) is mainly composed of Ti (53.98 at.%) and Si (36.67 at.%) with a atomic ratio of about 5:3. Therefore the phase can be identified as Ti5Si3. As shown in Figs.5 (b-f), its area fraction presents the rising tendency with increasing the Si content (11.30% for Coating II, 15.70% for Coating III, 20.84% for Coating IV, 29.84% for Coating V and 33.70% for Coating VI). Along with the change, TiB, TiB2 and TiNi are gradually reduced in area fraction, and finally completely disappear when the addition content of Si reaches 25 wt.%. However, the area fractions of Ti2Ni and TiC are approximately stable (about 65.54% for TiNi2 and 3.00% for TiC). The changes in area fraction of different phases are well in accordance with the XRD results. The evolution in phase constituent will

produce the essential effect on mechanical properties in terms of hardness and fracture toughness, and oxidation resistance as analyzed in the latter. 3.2 Mechanical properties The microhardness distribution throughout the whole cross sections of the laserclad samples is shown in Fig. 6. With respect the change in microhardness, the whole cross sections can be divided into two zones, namely the coating (1.80 mm), the transition zone (0.35 mm) and the substrate. The profile of the cross section can also be confirmed by the line scanning result of EDS. As shown in Fig. 7, the contents of Ni and Si are comparatively stable and fluctuate up and down around the certain values in the zone with a thickness of about 1.80 mm. When the distance from the surface of the coating exceeds 1.80 mm, their contents present a gradual downward trend. This trend is maintained at the zone with a thickness of 0.30 mm. Then their contents become very stable. With respect to Ti, its content is also comparatively stable in the zone (1.80 mm), then gradually increased in the zone (0.30 mm), finally maintains at a very stable value. The microhardness profile is well in accordance with the line scanning result of EDS. It is clear that the microhardness value fluctuates slightly across the whole cross sections of all coatings, indicating that the microstructure is uniformly distributed. Compared with the substrate (350 HV0.2), the average hardness of all coatings is significantly increased about double. The average hardness of Coatings I-VI is 960.55 HV0.2, 971.31 HV0.2, 980.96 HV0.2, 987.86 HV0.2, 1014.78 HV0.2, and 1031.42 HV0.2，respectively, which is about three times as much as that of the substrate. It is noted that the average hardness of the coatings has a slight increase with increasing the Si content. The average microhardness of all coatings can be calculated as 991.15 HV0.2. As far as the tested microhardness values are concerned, some statistical parameters (with respect to relative average deviation,

relative standard deviation) can be calculated as follows: 2.15% and 2.48%. This means that there is no significant change among the microhardness values of all coatings. The phenomenon may be associated with the change in phase constituent. As analyzed above, the contents of TiB2 and TiB with a very high hardness of about 2208.0 HV0.01 and 1287.2 HV0.01 [34] are gradually reduced with increasing the Si content, which is not beneficial to the improvement in hardness of the coatings. However, the content of Ti5Si3 with a comparatively high hardness of about 988.0 HV0.01 [35] is corresponding increased, which contributes to improving the microhardness of the coatings. Considering that the increase margin of the Ti5Si3 content is significantly higher than that of the TiB2 and TiB content, the synergistic effect causes the slight increase in microhardness of the coatings. The coatings’ fracture toughness (KIC) can be approximately evaluated by Vickers indentation method. The schematic diagram is shown in Fig. 8. A big load of N was applied to the different zones of the coatings’ cross sections on a Vickers hardness tester. The clear indentation can be prepared, and the cracks initiate and propagate around the four corners of the indentation due to the stress concentration. The KIC value can be calculated by the following formula [36]:  p K IC = 0.079 p  3  2 a

  log  4.5 a   c  

(9)

In which KIC represents the fracture toughness, p signifies the load; a denotes the half of the indentation’s diagonal length; and c is the half of the total length of the diagonal and the crack. In order to reduce the measuring errors and obtain the precise values, three indentations were measured on the cross section of each coating (shown in Fig. 9). The a values are approximately identical for all coatings. However, the c values

present the upward trend along with the increasing the Si content. Especially when the Si content is higher than 20 wt.%, the zone around the indentation apparently collapse and the length of cracks is also correspondingly prolonged, which imply the sharp drop in fracture toughness. The calculated KIC values also prove this. As demonstrated in Table 2, KIC of Coating I is about 4.788 MPa•m1/2, compared with which the value of Coating IV is reduced about 26.23%. However when the Si content is further enhanced to 30 wt.%, the value is sharply slowed down about 49.20%. The results indicate that the addition of Si causes the decrease in fracture toughness. Similarly, this tendency should be attributed to the relative change in content of TiB2 and Ti5Si3. Ti5Si3 with a complex hexagonal structure only has two slip surfaces of (1101) and (2311), and demonstrates the poor deformability in other directions. In addition, there is a strong covalent bond (Ti-Si) bond on the a-axis and a weaker metal bond (Ti-Ti) bond on the c-axis, causing a large non-harmonic motion in the c-direction. As a result, Ti5Si3 possesses high brittleness (a low KIC value of about 2.3 MPa•m1/2 [37]) and tends to cleavage fracture when it is subject to the external or internal force. Comparatively speaking, TiB2 demonstrates a higher KIC value of about 3.5 MPa•m1/2 [36] than Ti5Si3. The decrease in content of TiB2 can be considered as an important factor causing the reduction in fracture toughness of the coatings with the higher content of Si. Moreover, the decrease in content of TiNi is partially responsible for the change. It is well known that TiNi possesses the high plasticity and toughness due to its simple cubic structure with twenty slip systems. Less TiNi involved in the coating with a high content of Si addition results in the reduction in fracture toughness. Based on the above-mentioned analyses, it can be confirmed that the suitable Si content should be lower than 20 wt.%, in which the coating demonstrates the good comprehensive mechanical properties.

3.3 Oxidative resistance Fig.10 demonstrates the oxidation weight gain of the coatings in the isothermal oxidation tests carried out for 200 h at 800 °C. The oxidation weight gain of the coatings illustrates a regular change due to the addition of Si. The weight gain of Coating I is significantly higher than that of the other coatings at any time. Its final weight gain is about 7.49 mg•cm-2, which is about 1.3 times that of the other coatings. Considering that a new phase of Ti5Si3 is only synthesized in the other coatings, it implies that the introduction of Ti5Si3 can efficiently postpone the oxidation process. Accompanied with the synthesis of Ti5Si3, the further increase in Si content can still improve oxidation resistance of the coatings. The weight gain of the coatings is decreased from 5.79 mg•cm-2 to 4.72 mg•cm-2when the Si content is increased from 10 wt.% to 30 wt.%. However, it is worth noticing that the reduction in weight gain is not obvious when the addition content of Si exceeds 20 wt.% (only a slight reduction of 6.72% between Coating VI and Coating IV). Combined with the testing results of mechanical properties, the suitable addition content of Si should be determined to be 20 wt.%. The oxidation weight gain curve can be fitted well and expressed by the equations (Fig. 11). Two stages are involved in the whole oxidation process, referring to the initial oxidation stage and the slow oxidation stage. For the initial oxidation stage, the coating suffers from the serious oxidation due to the naked coatings strongly reacting with oxygen. Along with the oxides gradually formed, the coating will be covered with a thin layer of oxides, which will greatly retard the oxidation rate. The oxidation is correspondingly transformed into the slow oxidation stage. The curves can be described by the following two general equations: G = K1T in the initial oxidation stage

(10)

G 2 = K 2T in the slow oxidation stage

(11)

where G represents the oxidation weight gain, T is the oxidation time, K1 and K2 denote the constants of oxidation weight gain in the two stages. The specific fitting results with respect to every coating are indicated in Fig. 11. It is worth noted that the constant of T1 in the initial oxidation stage is approximately same with increasing the Si addition content from 5 wt.% to 30 wt.%, accompanied with which the duration is shortened from 22 h to 8 h. Moreover, the constant of K2 in the slow oxidation stage also demonstrates the reducing trend with the increasing in Si content. This implies that the addition of Si will play the positive role in improving oxidation resistance of the coatings. In order to reveal the oxidation mechanism, the coating subject to oxidation was examined by XPS. As shown in the survey spectrum (Fig.12 (a)), the oxides are composed of O, Ti, Al, Ni, Cr, and Si. The narrow spectra of above metal element were detected to confirm their chemical valence state (Figs.12 (b-f)), in which the black line refers to the experimental results and the red line represents the fitting results. For the Ti2p spectrum (Fig.12 (b)), two strong peaks located at 464.0 eV and 458.0 eV are clearly observed, which can be confirmed as two characteristic peaks of Ti in TiO2. A clear peak appears at 74.0 eV in the Al2p spectrum, confirming the existence of Al2O3 (Fig.12 (c)). With respect to the Ni2p spectrum (Fig.12 (d)), there are two peaks situated at 855.3 eV and 874.1 eV, indicating that the sample contains NiO. Similarly, two characteristic peaks related to Cr2p (577.3 eV) and Si2p (102.0 eV) imply that they are oxidized into Cr2O3 and SiO2 (Figs.12 (e) and (f)), respectively. The XPS results prove that the coating suffering from the oxidation is mainly composed of TiO2, Al2O3, NiO, Cr2O3 and SiO2. By comparing with the phase

constituents prior to the oxidation, it can be concluded that above-mentioned oxides are synthesized by the following reactions: TiNi+O2=TiO2+NiO

(12)

Ti2Ni+O2=TiO2+NiO

(13)

TiB+O2=TiO2+B2O3

(14)

TiB2+O2=TiO2+B2O3

(15)

TiC+O2=TiO2+CO2

(16)

Ti5Si3+O2=TiO2+SiO2

(17)

Fig.13 illustrates the relationship between change in standard Gibbs free energy (∆G0) and temperature (T) for the above reactions. From the point view of thermodynamics, all reaction can occur spontaneously over a wide temperature range due to their ∆G values less than zero. However, there is a significant difference in reaction priority at the given testing temperature (800 °C), namely Reaction (17)>Reaction (14)>Reaction (13)>Reaction (15)>Reaction (12)>Reaction (16). Obviously, Reaction (17) is easier to occur when compared with the other reactions, suggesting that Ti5Si3 can be regarded as the most unstable phase among all phases in the coating. This means that the coating with more Ti5Si3 may obtain the higher oxidation weight gain in the initial oxidation stage due to Ti5Si3 preferentially reacting with oxygen. As shown in Fig. 11, a shorter duration is obtained in a coating with a higher Si content in the initial oxidation stage. This implies that more addition of Si can accelerate the transformation of the coating from drastic oxidation to slow oxidation, and further retard the oxidation rate in the slow oxidation stage due to the decrease in the constant of K2 in the slow oxidation. The phenomenon is closely associated with the change in compactness of the oxide layer resulting from the

addition of Si. With the oxidation processing, a thin layer of oxides will be formed on the coatings’ surfaces, and separate the parts below the oxide layer from oxygen, which will greatly delay the invasion rate of oxygen into the coatings. The isolation effect is closely associated with the compactness of the oxide layer. Obviously, a layer of oxides with a higher compactness can better prevent the coatings from oxidizing. The compactness of the oxide layer can be characterized by a parameter named as PBR (Pilling-Bedworth ratio), which refers to the change in volume of a binary alloy before and after oxidizing. For a given binary alloy being subject to the oxidation (AxBy (s)+O2 (g)→AmOn (s)+BwOz (s)), the PBR value of the reaction system can be calculated by the following formula [38]: PBR Ax B y (s ) =

CV AmOn (s ) + DVBwOz (s ) EV Ax B y (s )

(18)

where VAx B y ( s ) , V AmOn (s ) and VBwOz ( s ) represent the molar volume of the solid reactant of AxBy, and the solid products of AmOn and BwOz, respectively. C, D and E signify the mole number of above-mentioned substances involved in the oxidation reaction.

VAx B y ( s ) , V AmOn (s ) and VBwOz ( s ) can be expressed as:

VAx B y (s ) =

M Ax B y (s )

ρA B x

VAmOn ( s ) =

M AmOn (s )

ρ A O (s ) m

VBwOz ( s ) =

(20)

n

M B wO z ( s )

ρ B O (s ) w

(19)

y (s )

z

(21)

in which M Ax B y ( s ) , M AmOn (s ) , M BwOz (s ) refer to the mole weight of those substances,

ρA B x

y

(s ) ,

ρ A O (s ) , ρ B O ( s ) are the density of those substances. m

n

w

z

Based on the above formulas, the PBR values related to or the systems in terms of Reactions (12)-(17) can be calculated of about 1.83, 1.90, 1.59, 1.34, 0.71 and 2.48. The value of Reaction (16) is less than 1, implying that the oxidation of TiC will cause the system shrunk in volume. This is to say that Reaction (16) will make the oxide layer become loose, result from which oxygen can still easily pass through the oxide layer and react with the inner of the coating. On the contrary, the other reactions with the values more than 1 will result in the volume expansion of the systems, which means that oxygen can be effectively isolated due to a compact oxide layer formed. Obviously, more Ti5Si3 involved in the coatings will promote the compactness of the oxide layers owing to its higher PBR value than those of the other reactions. This can well explain why oxidation resistance of the coatings is gradually improved along adding more Si content. However, it should be noticed that the growth compressive stress will generate in the oxide layer with a high PBR value during the oxidation process, which may cause the cracks initiate and propagate. As a result, oxygen may diffuse into the coating and accelerate the oxidation throughout the cracks. However, no cracks are observed throughout the whole cross sections of the coatings subject to the oxidation (shown in Fig. 14). Many fine pores are observed in the cross section of the oxidation layer in the coating with 5 wt.% Si (especially rich in the upper part), which may result from the volatilization of the formed gaseous oxides during the oxidation (such as B2O3, CO2). The pores provide the diffusion tunnels for oxygen, accelerating the oxidation of the coatings. The maximum thickness of the oxidation layer is about 55 µm. When the Si content is increased to 30 wt.%, the thickness of the oxidation layer is reduced to about 40 µm, and the pores are greatly increased in

number, which indicate that the compactness of the oxidation layer is significantly improved along increasing the Si content. That the coatings free of cracks may be attributed to the high breaking strength of the oxide layer. Moreover, the tensile stress resulting from Reaction (16) with a PBR value of less than 1 may counteract a portion of compression stress. Besides the above-mentioned reasons, the transformation of SiO2 from the crystalline state to the glassy state also play the positive role in inhibiting the generation of cracks due to the high liquidity and strong renewability of the glassy SiO2. The above analyses clearly show that oxidation resistance of the laser-clad coatings can be effectively improved by adding more Si (less than 30 wt.%) from the point of the volume change in initial phases before and after oxidation. However, considering that the high cracking susceptibility of the coatings with the Si content higher than 20 wt.%, the optimum content of Si can be confirmed as 20 wt.%.

4. Conclusions 1) The composites coatings were successfully fabricated on Ti6Al4V by laser cladding with a mixture of NiCrBSi and Si (5 wt.%-30 wt.%) as the cladding materials. Phase constituents of the coatings strongly depended on the addition content of Si. Ti2Ni, TiNi, TiB2, TiB and TiB were involved in the coatings with a Si addition content of 5 wt.%. A new phase of Ti5Si3 was synthesized when the Si addition content reached 10 wt.%, and the Ti5Si3 content presented a upward tendency with a decrease in content of TiB2, TiB and TiNi with increasing the Si content. 2) The average microhardness of the coatings was gradually improved along the increase in Si content (960.55 HV0.2 for Coating I, 971.31 HV0.2 for Coating II, 980.96 HV0.2 for Coating III, 987.86 HV0.2 for Coating IV, 1014.78 HV0.2 for Coating V, and 1031.42 HV0.2 for Coating VI), coupled with the increase in fracture toughness ( 4.788 MPa•m1/2 for Coating I, 4.772 MPa•m1/2 for Coating II, 4.760 MPa•m1/2 for

Coating III, 3.532 MPa•m1/2 for Coating IV, 2.720 MPa•m1/2 for Coating V, and 2.432 MPa•m1/2 for Coating VI). 3) Oxidation resistance of the coatings was closely associated with the addition content of Si. The increase in Si content accelerated the formation of an oxide layer on the coatings’ surface, resulting in the oxidation rapidly entering into the slow oxidation stage with a lower constant of oxidation weight gain. The oxidation resistance of the coatings was gradually improved with increasing the Si content due to the oxidation weight gain reduced from 7.49 mg•cm-2 to 4.72 mg•cm-2 with the Si content increased from 5 wt.% to 30 wt.%. A compacter oxidation layer was formed on the surface of the coating with a higher Si content due to the higher PBR value for Ti5Si3 subject to oxidation than those of the other phases, which played the essential role in the improvement in oxidation resistance of the coatings. 4) The suitable content of Si was confirmed as 20 wt.%, in which the excellent comprehensive mechanical properties and outstanding oxidation resistance were obtained.

Acknowledgments This work is financially supported by the National Natural Science Foundation of China (51471105) and “Shu Guang” project of Shanghai Municipal Education Commission and Shanghai Education Development Foundation (12SG44).

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Figure captions: Fig.1. X-ray diffraction results of the coatings. Fig.2. The relationship between change in Standard Gibbs free energy of Reaction (1)-(8) and temperature. Fig.3. Morphologies of the cross section of Coating I: (a) cross sectional view of the sample; (b) micrograph of the interface between the coating and the substrate. Fig.4. Microstructure of the coatings’ cross sections: (a) Coating I; (b) Coating II; (c) Coating III; (d) Coating IV; (e) Coating V; (f) Coating VI. Fig.5. The evolution diagram of the area fraction of different phases. Fig.6. Microhardness distribution across the cross sections of all coatings. Fig.7. The line scanning result of EDS for Coating I: (a) cross sectional view of the sample; (b) the distribution of Ti, Ni and Si across the whole cross section of the sample. Fig.8. Schematic diagram of the fracture toughness (KIC) test. Fig.9. The indentation optical images:(a) 5 wt.% Si；(b) 10 wt.% Si；(c) 15 wt.% Si ；(d) 20 wt.% Si；(e) 25 wt.% Si；(f) 30 wt.% Si. Fig.10. Oxidation weight gain of the coatings in isothermal oxidation tests at 800°C. Fig.11. Fitting curves of oxidation weight gain and oxidation time for different samples exposed at 800 °C for 200 h: (a) 5 wt.% Si；(b) 10 wt.% Si；(c) 15 wt.% Si ；(d) 20 wt.% Si；(e) 25 wt.% Si；(f) 30 wt.% Si. Fig.12. XPS spectra of oxide films formed on the coatings’ surfaces (a) Full spectrum ；(b) Ti2p；(c) Al2p；(d) Ni2p；(e) Cr2p；(f) Si2p. Fig.13. The relationship between change in Standard Gibbs free energy of Reaction (12)-(17) and temperature.

Fig.14. The morphologies of the coating’s cross sections subject to the oxidation: (a) 5 wt% Si; (b) 30 wt.% Si.

Tables Table 1 EDS results for the different zone marked in Figs.3 (a2) and (b2) (in at. %). Regions

Ti

B

C

Ni

Al

Si

V

Cr

1

41.52

—

4.89

24.60

8.96

5.63

3.05

11.34

2

31.69

—

2.34

47.71

14.41

0.94

0.61

2.31

3

74.81

—

3.38

9.39

3.51

3.28

1.37

4.26

4

9.89

88.13

1.22

0.19

0.03

0.10

0.31

0.13

5

10.78

84.97

1.11

0.35

0.07

0.19

1.13

1.39

6

53.98

—

2.15

1.90

1.71

36.67

2.31

1.28

Table 2 Measurement results of KIC values for different coatings KIC/MPa•m1/2

Coatings

Average value

I

4.620

5.300

4.445

4.788

II

5.102

4.89

4.325

4.772

III

4.979

4.886

4.417

4.760

IV

3.321

3.350

3.925

3.532

V

2.431

2.577

3.152

2.720

VI

2.241

2.593

2.464

2.432

Highlights A laser-clad coating with outstanding oxidation resistance was produced on Ti6Al4V. The coating also exhibited good mechanical property (hardness and toughness). Microstructural evolution with increasing Si content in cladding material was explored. Changes in mechanical property resulting from microstructural evolution were studied. The oxidation mechanism of the coatings with the Si content was revealed in details.

Declaration of Interest Statement All authors have read and approved this version of the article, and due care has been taken to ensure the integrity of the work. Neither the entire article nor any part of its content has been published or has been accepted elsewhere. It is not being submitted to any other journal. We believe the paper may be of particular interest to the readers of your journal. Jun Li