Surface & Coatings Technology 291 (2016) 43–53

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Surface & Coatings Technology

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Formationmechanism andmicrostructural and mechanical properties ofin-situ Ti–Ni-based composite coatings by laser metal deposition

Chenglong Ma a,b, Dongdong Gu a,b,⁎, Chen Hong c, Beibei He a,b, Kun Chang a,b, Qimin Shi a,b

a College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing 210016, PR Chinab Institute of Additive Manufacturing (3D Printing), Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing 210016, PR Chinac Fraunhofer Institute for Laser Technology ILT/Chair for Laser Technology LLT, RWTH Aachen, Steinbachstraße 15, D-52074 Aachen, Germany

⁎ Corresponding author at: College of Materials ScieUniversity of Aeronautics and Astronautics, Yudao Street

E-mail address: dongdonggu@nuaa.edu.cn (D. Gu).

http://dx.doi.org/10.1016/j.surfcoat.2016.02.0130257-8972/© 2016 Elsevier B.V. All rights reserved.

a b s t r a c t

a r t i c l e i n f o

Article history:Received 20 December 2015Revised 26 January 2016Accepted in revised form 5 February 2016Available online 6 February 2016

The in-situ TiO2 reinforced Ti–Ni composite coating on carbon steel was successfully prepared by laser metaldeposition (LMD) using Ti–Ni as-mixed powder with an atomic ratio of 60:40. With the aim of in-situ reactiondesign during LMD processing, a trace of oxygen mixed with the shielding gas was introduced. Different “laserenergy input per unit length” (E) by changing the laser power was set to investigate the influence on the depo-sition quality and attendant microstructure and mechanical property of the LMD-processed layer. TiO2 particleswith unique flower-like structure formed when the applied laser energy E ≤ 96 kJ/m, while the apparentoxidation of grain boundaries was observed as E increased to 120 kJ/m. The formation mechanism of in-situTiO2 particles with a flower-like structure was present. At the optimized process parameter of 96 kJ/m, theLMD-processed layer showed the highest densification degree free of any pores and cracks. The correspondingmechanical properties were measured, showing the relatively high average microhardness of HV0.2 790 andsignificantly improved tribological property containing lower coefficient of friction of 0.4 and more smoothworn surface.

© 2016 Elsevier B.V. All rights reserved.

Keywords:Laser metal deposition (LMD)Ti–Ni alloyCoatingComposite

1. Introduction

Ferrous alloys, as the typical structuralmaterial, have gotwide appli-cations in engineering field due to their excellent comprehensive prop-erties and low cost. Normally, the components made of ferrous alloysare exposed to various kinds of external environments, which maylead to an early failure because of some reasons such as corrosion andabrasion. Twomain approaches are generally used to improve the prop-erties of ferrous alloys, namely alloying and surfacemodification. Takinginto account of the cost, the surface modification has been paid moreattentions into [1–3].

Laser metal deposition (LMD), as one of themost prevailing additivemanufacturing (AM) technologies, has got rapid development in recentyears. Due to its highly versatile process capability, LMD has beenwidely applied to manufacture new components, to repair and rebuildworn or damaged components, and to prepare wear- and corrosion-resistant coatings [4–6]. When used for material surface property mod-ification, LMD is also known as laser cladding or laser depositing. Basedon the processing features of AM technology, namely layer-to-layer

nce and Technology, Nanjing29, Nanjing 210016, PR China.

shaping and consolidation of feedstock using a computer-controlledlaser as the energy source, LMD possesses the nature of highly non-equilibrium, rapid melting and solidification [7–9]. Consequently,uniform distribution of chemical components and novel microstruc-tures, which are hard to be achieved by other conventional methods,can be introduced by LMD. Relative investigations on different typesof laser-deposited coatings or layers on the surface of ferrous alloyshave been reported. For example, Qunshuang Ma et al. prepared theNi60/WC composite coatings on the surface of Q550 steel. A specialcored-eutectic structure was obtained, which had advantages of well-distribution and tight binding with the matrix [10]. Dariusz Bartkowskiet al. successfully achieved laser cladding Stellite-6/WC compositecoatings on the surface of low-carbon steel. This study indicated thatthe microhardness enhanced with the volume fraction of WC increaseswhile the corrosion resistance weakened [11].

Titanium and its alloys, due to the excellent specific strength, frac-ture resistance characteristics and high corrosion resistance, have wonextensive applications in engineering field [12,13]. However, the poortribological properties including high friction coefficient, low adhesiveand frettingwear resistance restricted the application of Titanium alloysas coatings to a large extent [14,15]. Ti–Ni alloys have been proven to bean alternative coating material, owing to the significant intensificationof tribological properties by the formation of Ti–Ni intermetallic. Thecorrelative research works have been carried out. M. Salehi et al.successfully prepared Ti–Ni intermetallic coatings on the surface of a

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carbon tool steel by a duplex process. During the duplex process,the Ni coating with the thickness of 20 μm was firstly achieved byelectroplating, and then Ni-coated specimens were packed in the Ti-rich powder mixtures in order to develop Ti–Ni intermetallic coatings.Besides, diffusion annealing was carried out to reduce compositionalgradient profile from surface to core and the final phases were com-posed of TiNi, Ti2Ni, TiNi3 and FeNi [16]. F. T. Cheng et al. systematicallyinvestigated laser cladding of AISI 316 stainless steel using respectivelypreplaced NiTi wire, NiTi strips and NiTi powder. In above three condi-tions, a good coating with high microhardness and tough interface, freeof any pores and cracks, was all obtained. As a result, the surface cavita-tion erosion resistance of AISI 316 stainless steel achieved a significantimprovement [17–19]. Further improvement has been obtained by theaddition of the third alloying element or reinforcement particles, suchas N, Al and ZrO2. Cunshan Wang et al. reasonably designed theTi70.3Ni22.2Al7.5 alloy, optimized from a basic binary eutectic Ti76Ni24alloyed with different amounts of Al, and prepared the correspondingcoating on AZ91HP magnesium alloy by laser cladding. The coatingmainly consists of β-Ti solid solution and Ti2Ni intermetallic compoundresulting in high hardness, good wear resistance and corrosion resis-tance [20]. Babatunde A. Obadele et al. studied the influence of laserscan speed andadditive ZrO2 particle onmicrostructure,microhardness,wear resistance and corrosion resistance for laser deposited Ti–Nicoating on the surface of AISI 316 stainless steel. When the laser scanspeed reached 0.0067m/s, the Ti–Ni coating showed themost excellentcorrosion resistance and the addition of ZrO2 was able to enhancedramatically the corrosion resistance [14].

According to the above investigations, the addition of the thirdcomponent in Ti–Ni alloy can introduce prominent “gain effect” andefficiently strengthen the mechanical or chemical property of Ti–Nialloy. However, the investigations on the Ti–Ni composite coatingreinforced by particles are still lacking, especially on non-equiatomicTi–Ni composite coating as well as for the application on surface modi-fication of ferrous alloy. In this study, in-situ TiO2 reinforced Ti–Nicomposite coating on carbon steel was successfully prepared by LMDusing Ti–Ni as-mixed powder. By the introduction of little oxygenmixed with the shielding gas, in-situ TiO2 particles with differentstructures were observed to form within the matrix during the LMDprocessing with different process parameters, which had an importanteffect on the deposition quality and attendant mechanical property ofthe laser-deposited layer. TiO2 particle, as a reinforcing phase, hasbeen applied for some other metal matrixes, such as A356 aluminumalloy [21], but few applications for Ti–Ni alloy matrix have been report-ed. Furthermore, this result will be helpful to prepare in-situ oxideparticles reinforced Ti–Ni composite coating by the introduction of ap-propriate oxygen carrier during laser processing. Hence, in this presentpaper, different laser processing parameters were set in order to inves-tigate the influence on the deposition quality and attendantmicrostruc-ture and mechanical property of the deposited layer. Meanwhile, theformation mechanism of in-situ TiO2 particles with a flower-like struc-ture was present.

2. Experiment procedure

2.1. Materials

The gas-atomized, spherical Titanium powder (99.5% purity) with amean particle size of 30 μm and the spherical-shaped Nickel powder(99.5% purity) in an average particle size of 10 μm were used asfeedstock powders in this study. Titanium powder and Nickel powderin an atomic ratio of 60:40 were uniformly mixed in a high-energyPulverisette 6 planetary mono-mill (Fritsch GmbH, Germany) using aball-to-powder weight ratio of 5:1, a ration speed of the main disk of200 r/min, and a milling time of 4 h. The as-mixed powder stilldemonstrated original spherical morphology, as shown in Fig. 1b.Besides, carbon steel was utilized as the substrate material, which was

sandblasted and ultrasonically cleaned successively in acetone and de-ionizedwater to remove surface contaminants before LMD. The detailedchemical compositions of raw powders and substrate were present inTable 1.

2.2. Laser metal deposition

The schematic of the experimental procedure for laser depositing Ti–Ni alloy layer on the surface of carbon steel substrate was present inFig. 1a. Laser metal deposition (LMD) processing was carried out witha 5-axis CNC system, a Trumpf Nd:YAG laser system with a maximumoutput power of 3 kW and a focused spot diameter of 0.6 mm, integrat-ed with a powder feeder system, and a coaxial powder nozzle. Thepremixed Ti–Ni powder was injected into the melted pool throughthe coaxial nozzle with a powder feeding rate of 2.4 g/min. At thesame time, with the aim of in-situ reaction design, Argon mixed witha trace of oxygen acted as carrier gas during laser deposition. Themultiple tracks were cladded for the layer with the dimension of5 mm × 28 mm. LMD fabrication was based on the line-by-line lasercladding, using a constant laser scan speed of 500 mm/min and variouslaser powers of 400W, 600W, 800W, 1000Wand1200W. An integrat-ed parameter “laser energy input per unit length” (E), which wasdefined by E = P/ν, as shown in Table 2, was used to estimate thelaser energy input to the deposited layer. Although five samples withdifferent processing parameters were prepared in this study, themicrostructure features of sample N1 with 48 kJ/m were similar assample N2 with 72 kJ/m and the microstructure features of sample N4with 120 kJ/m were similar as sample N5 with 144 kJ/m. Therefore,three samples (N2, N3, and N4), as the typical ones, were chosenfor the microstructure study. The macrophotograph of final LMD-fabricated layers was present in Fig. 1c.

2.3. Characterization of microstructure, chemical composition and phase

The phases of the deposited layer were identified by a Bruker D8Advance X-ray diffractometer (XRD) with Cu Kα radiation (λ = 0.15418 nm) at 40 kV and 40 mA using a continuous scan mode. Threespecimens (N1, N2, and N3) for metallographic examinations wereprepared according to the standard procedures and etched with asolution composing HF, HNO3, and distilled water with a volume ratioof 1:6:7 for 3 s. A PMG3 optical microscope (Olympus Corporation,Japan) was used to observe the low-magnification three-dimensionmorphology and cross-sectional microstructures of LMD-processedlayer on the carbon steel. High-resolution studies of themicrostructuralfeatures of the deposited layer were characterized using an S-4800 fieldemission scanning electron microscope (FE-SEM) (Hitachi, Japan) at anaccelerating voltage of 5.0 kV. Chemical compositions were determinedby an EDAX's energy dispersive X-ray spectroscope (EDX) (EDAX, Inc.,USA).

2.4. Microhardness and dry sliding wear test

The Vickers hardness was measured using a MicroMet 5101 micro-hardness tester (Buehler GmbH, Germany) at a load of 0.2 kg and anindentation time of 15 s. The tribological property of the specimenwas estimated by the dry sliding wear tests conducted in a HT-500ball-on-disk tribometer (Lanzhou ZhongKe KaiHua Sci. &Technol. Co.,Ltd., China) in air at room temperature. The counter material wasGCr15 bearing steel ball with a diameter of 3 mm and a mean hardnessof HRC 60, using a test load of 220 g was applied. The friction unit wasrotated at a speed of 560 rpm for 15 min, with the rotation radius of2 mm. The coefficient of friction (COF) of the specimens was recordedduring wear tests.

Fig. 1. (a) The schematic diagram of LMD processing Ti–Ni layer; (b) the raw as-mixed Ti/Ni powder with an atomic ratio of 60:40; (c) the final LMD-processed Ti–Ni layers.

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3. Results and discussion

3.1. The quality of deposition and element distribution

Figs. 2a–4a demonstrate the three-dimension (3D) morphology oflaser deposited layer at laser energy input per unit length (E) of 72, 96and 120 kJ/m, respectively. Furthermore, to better reflect the qualityof deposition, the corresponding composition profiles along the depthof the deposited layer as determined by EDX are also shown inFigs. 2b–4b. The parabolic shape bond line between the depositedlayer and the substrate, characteristic microstructure feature of laserprocessing, was clearly observed from the cross-section of each sample(Figs. 2–4a). Although adhesion of the deposited layer and the substratewas good at each parameter E, the overall quality of the deposited layervaried.When the applied Ewas 72 kJ/m, some cracks propagating in theoverlap zone between two adjacent tracks could be obviously seen(Fig. 2a). Besides, some agminated pores were found at the beginningposition of the crack lying in X-direction cross section. Then the EDXdetection indicated that the concentrations of Fe, Ti, and Ni elementsalong the depth of the deposited layer presented an apparent, near-invariable gradient (Fig. 2b). As E further increased to 96 kJ/m, a nearlyfull dense layer was achieved, free of any cracks and pores (Fig. 3a). Inthis condition, a transfer zone with an approximately 300 μm depthshowed up and the concentrations of all elements tended to be invari-able when the distance from the interface was over 300 μm (Fig. 3b).However, cracks running through the deposited layer were observedat the Y-direction cross section when the applied E reached 120 kJ/m(Fig. 4a). In this case a more homogenous element distribution withhigh Fe content was obtained (Fig. 4b).

During LMD processing, both part of substrate material and coaxialfeeding powder were melted, forming a molten pool with a continuousliquid front. The different laser processing parameter determined differ-ent amount of the liquid formed in themoltenpool and resultant quality

Table 1The chemical compositions of raw powders and substrate.

Composition in wt.%

Fe C Si Mn Cr Ni Cu Ti

Carbon steel Balance 0.88 0.17 0.6 0.12 0.1 0.25 0Ni–Ti mixed powder 0 0 0 0 0 55.03 0 44.97

of deposition. When the E was relatively low, the amount of the liquidwithin the molten pool was limited due to insufficient laser energyinput, which made the liquid hardly spread out smoothly. As a result,the densification behavior within the molten pool weakened alongwith the emergence of pores and meanwhile the homogenization rateof all elements was restricted dramatically. Moreover, the fierce shrink-age induced by the melt solidifying rapidly along the horizontal andvertical direction of themolten pool led to the formation of tensile stressin the deposited layer. Specially, the overlap zone usually experiencedthe remelting and solidification process repeatedly, thus causing ahigher stress concentration [22]. Then the combined effect of poresand tensile stress led to the initiation and propagation of cracks. Onthe other hand, when an excessive E was applied, more laser energywas absorbed by the substrate material and powder and then the tem-peraturewithin themolten pool rose rapidly. Consequently, the formedmelt was in a state of overheating, thus leading to the accumulation ofthermal stress during subsequent solidification processing. On themax-imum stress accumulated in the solidified layer exceeded the strengthof the deposited alloy, cracks occurred immediately with the release ofthe accumulated stress. In this case, the excessive heat input acceleratedvastly the diffusion of elements, thus achieving a more homogenousdistribution state of elements.

Besides, the height of the deposited layers shows an apparent fluctu-ation in Figs.2–4. This fluctuation is mainly attributed to the combinedeffect of surface tension and powder vaporization [23]. Due to the intro-duction of oxygen in this study, the surface tension gradient may showpositive value, thus leading to formation of theMarangoniflow from theedge of melt pool to the center of melt pool and an increase of height ofmelt pool. With the laser power increasing, surface tension gradientenhances significantly and therefore the increasing effect of height ofmelt pool is more apparent. However, when excessive laser energy is

Table 2The LMD process parameter.

Sample Laser power (W) Scan speed (mm/min) Energy density (kJ/m)

N1 400 500 48N2 600 500 72N3 800 500 96N4 1000 500 120N5 1200 500 144

Fig. 2. (a) 3-D optical metallograph (OM image) composite showing the deposition quality of Ti–Ni layer at the relatively low E of 72 kJ/m; (b) the corresponding composition profilesalong the depth of the deposited layer.

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input, the powder vaporization effect gets intensified remarkably andconsequently decreases the height of melt pool.

3.2. Microstructure evolution along the depth of the deposited layer

According to the above analysis, at an optimized E of 96 kJ/m, the op-tical micrograph of the cross-section microstructure of the depositedlayer, locating at different regions from the surface, is shown in Fig. 5.In fact, the sample N2 has the similar cross-sectional microstructurefeatures as the sample N3. Grain morphology transformations of thedeposited layer from the top to the bottom are relatively coarseningsteering dendrites, directional equiaxed grains, fine dendrites andplanar crystals. The previous investigations have proved that thetemperature gradients (G) and the solidification rates (R) during

Fig. 3. (a) 3-D optical metallograph (OM image) composite showing the deposition quality of Tdepth of the deposited layer.

solidification can determine microstructure developments [24,25]. Thetemperature gradient G and the solidification rate R can be calculatedin the following equations [26]:

G ¼ 2K T−T0ð Þ2ηP

ð1Þ

R ¼ Vs cosθ ð2Þ

where K is the thermal conductivity of the material (W m−1 K−1), T isthe liquidus temperature of the alloy (K), T0 is the initial temperatureof the substrate (K), η is the laser absorption coefficient, P is the laserpower (W), Vs is the laser scanning speed (m s−1) and θ is the anglebetween Vs and R. In general, the value of G/R was used to evaluate

i–Ni layer at the proper E of 96 kJ/m; (b) the corresponding composition profiles along the

Fig. 4. (a) 3-D optical metallograph (OM image) composite showing the deposition quality of Ti–Ni layer at an excessive E of 120 kJ/m; (b) the corresponding composition profiles alongthe depth of the deposited layer.

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the influence on the morphology of the liquid/solid interface or micro-structure developments [26]. According to the above equations, G/Rcould be estimated by:

G=R ¼ 2K T−T0ð Þ2ηPVs cosθ

: ð3Þ

At the bottom of the molten pool (~50 μm from the interface), thesolidification rate R was nearly perpendicular to the laser scanningspeed Vs and the temperature gradient G reached the maximum valuedue to the contacting substrate, resultantly G/R getting the maximumvalue. Thus the melt was in a state of overheating and consequently

Fig. 5. The opticalmicrograph of the cross-sectionmicrostructure of the deposited layer at (a) thsurface, and (d) the bottom of the deposited layer under the proper E of 96 kJ/m.

the solidification microstructure was planar crystal. With the solidifica-tion front of the melt advancing (~300 μm from the interface), the pre-viously solidified layer acted as the substrate and consequently theinitial temperature T0 increased, thus leading to a decrease of tempera-ture gradientG located in the solidification front of themelt. At the sametime, the angle θ between Vs and R became lower, causing the increaseof solidification rate R. In this condition, the value of G/R decreased sig-nificantly and as a result constitutional undercooling occurred. In thiscase, dendrites growth pattern took over a leading position. When theliquid/solid interface continued to move, both the bottom and the topof themolten pool were being solidified due to the contacting substrateand the air, which led to the change of G direction and subsequently the

e near-surface region, (b) 1000 μmdistance from the surface, (c) 500 μmdistance from the

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formation of steering dendrites at the top of the molten pool (~50 μmfrom the surface). Then the solidification rate from the bottom to thetop was faster than that from the top to the bottom, leading to two sim-ilar temperature gradients in the upper part of the deposited layer(~200 μm from the surface) and resultant directional equiaxed grains[24].

3.3. Microstructure evolution and phase identification as E increasing

Also, to depict the influence of E by changing laser power on thecross-section microstructures of the deposited layer, the corre-sponding FE-SEM characterization is performed at the same distance(~300 μm) from the interface, as shown in Figs. 6–8. At a relativelylow E of 72 kJ/m, the main microstructure demonstrates typical den-drite and at the same time some dispersive particles with submicronscale are observed (Fig. 6a and b). High- magnification FE-SEMmicrograph shows the detailed microstructure of these dispersive

Fig. 6. (a) The cross-sectional microstructure of the deposited layer at ~300 μm distance from twas applied; (b) the high-magnification FE-SEMmicrograph of (a); (c) particlemorphologywitcharacterization of point 1 located in (c).

particles, a cluster structure by several finer particles getting togeth-er (Fig. 6c). The XRD experiment identifies Fe2Ti, FeNi3 and TiO2 asthe main constituent of the microstructure (Fig. 6d). Besides, a fewTi–Ni intermetallic phases, such as Ni4Ti3 and Ni3Ti, were also detect-ed. The elevated E (reaching 96 kJ/m) results in the transformationfrom typical dendrite to cellular dendrite (Fig. 7a and b). Interesting-ly, particles with unique flower-like structure are found to distributewithin the intercellular regions homogenously (Fig. 7c). The detect-ed main phases are still Fe2Ti, FeNi3 and TiO2 and meanwhile Ni4Ti3phase are also identified clearly, according to the XRD resultspresented in Fig. 7d. As an excessive E is applied (120 kJ/m), thesolidification microstructure is completely different, showing rela-tively loose microstructure (Fig. 8a). In this case, the diffraction in-tensity of Ti–Ni intermetallic phase is weak, as shown in Fig. 8b.Besides, the intergranular region was observed to get an apparentoxidation reaction, as shown in Fig. 8c, which was proved by thefurther EDX characterization (Fig. 8d).

he interface, showing the typical dendritic morphology when a relatively low E of 72 kJ/mh a cluster structure; (d) XRD patterns obtained over awide range of 2θ (20–90°); (e) EDX

Fig. 7. (a) The cross-sectional microstructure of the deposited layer at ~300 μm distance from the interface, showing the cellular dendrite morphology when a proper E of 96 kJ/m wasapplied; (b) the high-magnification FE-SEM micrograph of (a); (c) particle morphology with a flower-like structure; (d) XRD patterns obtained over a wide range of 2θ (20–90°).

Fig. 8. (a) The cross-sectional microstructure of the deposited layer at ~300 μm distance from the interface, showing the iron filing shaped morphology when an excessive E of 120 kJ/mwas applied; (b) the high-magnification SEM image of (a); (c) EDX characterization of point 2 located in (b); (d) XRD patterns obtained over a wide range of 2θ (20–90°).

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Based on the above transformation feature of grainmorphologywithE increasing, the typical dendrites development in sample N2 happensprior to that in N3 sample. The previous analysis has depicted theimportance of G and R on the grain morphology. Moreover, the solutioncomposition also influences the grain morphology [27]. The criticaltemperature gradient Gc was defined as [28]:

Gc ¼ −mlCS 1−kð Þ

kRDlð4Þ

wheremlwas the slop of the liquidus line, Clwas the liquid compositionin the liquid/solid interface, Cs = k Cl, the solid composition in theliquid/solid interface and Dl represented the diffusion coefficient. Con-stitutional undercooling occurred based on the premise of G b Gc meet-ing constitutional undercooling [29], thus leading to the development ofcellular or dendritic grains (the lower constitutional undercooling tendsto induce cellular grains while the higher constitutional undercoolingtends to induce dendritic grains). Due to the insufficient laser energyinput in sample N2, the maximum temperature within the moltenpool and attendant Marangoni convention is limited, which causesthe decrease of diffusion coefficient. Hence, the critical temperature gra-dient Gc is improved and consequently the constitutional undercoolingregion gets extended in N2 sample, which can account for the differencein the microstructures of N2 and N3 samples. However, it makes adifference when E reaches 120 kJ/m. As an excessive energy input isapplied, a large amount of substrate material is melt and the melt is ina state of overheating, which accelerates the oxidation behavior withinthe molten pool. A number of oxygen atoms aggregate in the grainboundaries to cause the oxidation of grain boundaries, which may bebelieved to account for the formation of the observed solidificationmicrostructure.

3.4. Formation mechanism of dispersive TiO2 submicronflower

With the aim of in-situ reaction design during the LMD process-ing, a trace of oxygen by controlling the airflow velocity of shieldinggas was introduced. EDX detection found that a number of particlesobserved in Figs. 6–8 belonged to an oxide phase TiO2 (Fig. 6e). The

Fig. 9. (a) The crystal structure of rutile; (b) the observed crystal morphology schematic o

corresponding XRD characterization had indicated that the oxidephase was rutile TiO2 with relative strong diffraction peaks (110) and(211) (Figs. 6d, 7d and 8b). Specially, TiO2 particles with the flower-like structure could be observed when the applied laser energyE ≤ 96 kJ/m (Figs. 6c and 7c). The flower-like structure became invisiblewhen an excessive E value of 120 kJ/m was applied. Fig. 9a shows thecrystal structure of rutile, a tetragon with Ti4+ surrounded by six O2−

at the corners of a slightly distorted octahedron and each O2−

surrounded by three Ti4+ lying in a plane at the corners of an equilateraltriangle [29]. As the molten pool formed by interaction between thelaser and the powder, dissociative Ti atoms tended to unitewith oxygenatoms, taking into account of the higher activity of Ti atom. The forma-tion processing of TiO2 might experience two stages:

Tiþ O→TiO ð5Þ

TiOþ O→TiO2: ð6Þ

TiO2, as a highmelting point phase, preferentially precipitated in themelt pool to form nanocrystalline. The observed crystal morphologyschematic of rutile is shown in Fig. 9b [30]. Due to the fact that the(110) crystal plane is the plane with the least energy, the growth rateof the (110) crystal plane along the b001N direction is higher thanthat of other planes [31]. Hence, the zone size of the (110) crystalplane was larger than that of other planes in Fig. 9b. In XRD detection,the main peak of TiO2 corresponding to the (110) crystal plane demon-strated the strongest intensity, thereby proving the above point. Inorder to reduce the number of crystal planes with high energy, a clusterstructure formed as more Ti and O atoms continued to deposit onthe (110) crystal planes, thus improving the stability of precipitatedTiO2 nanocrystalline. When the applied laser energy was elevated, thegrowth of the (110) crystal plane along the b001N direction got en-hanced due to the sufficient heat input during the LMD processing,thereby leading to the formation of submicronflower structure. Thecorresponding schematic diagram is shown in Fig. 9c. However, whenexcessive laser energy is applied, a mass of heat input leads to theintroduction of excessive Fe component and attendant improvementof activity of Fe atom. As a result, the combination of Fe and O atom is

f rutile; (c) the corresponding schematic diagram of particle morphology evolution.

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enhanced, which accounts for the formation of iron oxide in the inter-granular region. Besides, the Marangoni convention [32] induced bythe surface tension gradientwithin themolten pool was also intensifiedsignificantly, thus disturbing the direction of heat flow and hinderingthe formation of TiO2 submicronflower.

3.5. Mechanical property characterization

Fig. 10 depicts the microhardness distributions along the depth ofLMD-processed layers. For all laser-deposited samples, the averagemicrohardness achieved significant improvement compared with thesubstrate material showing a very low average value of 280 HV0.2. Theformation of Fe–Ti and Fe–Ni intermetallic as well as dispersive oxideparticles were the main factors for the improvement. The maximummicrohardness in N2 sample was measured on the near surface,reaching a level as high as 840.8 HV0.2, which was higher than that inother samples, implying that more Ti–Ni intermetallic phases mightform in the near surface of N2 sample in view of the XRD results. Forthe chosen typical processing parameters, the N4 sample had the lowestaverage microhardness in comparison with N2 and N3 samples. Due tothe significantly intensified oxidation behavior and resultant loosesolidification microstructure induced by an excessive laser energyinput, the capacity of plastic deformation resistance gets weakened.

Fig. 11 shows the coefficients of friction (COFs) as well as thecorresponding worn morphologies of the deposited layers and thesubstrate material. For the substrate material, the worn surface showeda very high roughnesswith severematerial delamination and spallation.As a result, relatively high COF of 0.8 was obtained. An apparentimprovement in tribological property was achieved when Ti–Ni alloycoating was prepared on the surface of substrate material. At the initialrunning-in stage of the wear test, the COF curves all exhibited fluctua-tion to some degree, which was attributed to the relatively roughsurface and the existence of oxide film [33]. With the sliding timeincreasing, the fluctuation of COF tended to be steady because of thedecreased surface roughness and the stripping of oxide film.

The worn surface morphologies of samples for typical processingparameters are shown in Fig 11. Nevertheless, the differences in COFvalues and worn morphologies indicated that the tribological propertyof the deposited layer was influenced significantly by the laser process-ing parameter. At the relatively low E of 72 kJ/m, the average COF valuereached a high level of 0.58 and a number of wear debris composed ofultrafine particles produced on the worn surface. As the E increased to96 kJ/m, the average COF value significantly decreased to 0.4, and

Fig. 10. The microhardness distributions along the depth of LMD-processed layer,respectively corresponding to N1, N2, N3, N4 and N5 sample.

the much smoother worn surface with shallow grooves was obtained,nearly free of any wear fragments. When the applied E increased to120 kJ/m, the average COF value increased to the higher level of 0.65and more wear debris as well as crumbling were found on the wornsurface.

The tribological property of the deposited layer is closely relatedto its densification level, microstructures and microhardness [33].When the E value is low, densification behavior of the LMD-processedlayer is restricted due to the formation and expanding of stress/pore-induced cracks. This leads to a remarkable increase of the surface rough-ness, thus enlarging the amount of abrasive fragments. As a result, morehard and brittle fine particles produce as the surface levels off and theCOF value shows an apparent fluctuation, showing a typical abrasivewear. Taking into account of the highest surface microhardness, furtherwear of the deposited layer is limited after the surface becomes relative-ly smooth. As the proper E is applied, sufficient laser energy guaranteesthe formation of good densification and homogenous microstructures,therefore leading to the enhanced wear resistance. At an excessivelaser energy input, densification level decreases again due to theexistence of cracks. Besides, the loose solidification microstructurewith oxidized grain boundaries as well as the lowest microhardnesscan be responsible for the relatively poor worn surface. However, thehomogenously dispersed oxide particles with submicrometer scale inN2 and N3 samples are believed to play a role not to be ignored inwear test. These oxide particles can improve efficiently the strength ofthe deposited layer and then enhance the capability of plastic deforma-tion resistance.

4. Conclusion

The in-situ TiO2 reinforced Ti–Ni composite coating on carbon steelwas successfully prepared by LMD using Ti–Ni as-mixed powder withan atomic ratio of 60:40. Different “laser energy input per unit length”(E) by changing the laser power was set to investigate the influenceon the deposition quality and attendantmicrostructure andmechanicalproperty of the LMD-processed layer. The main results can be summa-rized as:

(1) The highest densification level and homogenous element distri-bution were obtained when a proper E of 96 kJ/m was applied,while the application of relatively lower or higher E resulted inthe formation of cracks. Besides, the diffusion effect of Fe aswell as oxidation behavior in the deposited layer got enhancedas an excessive laser energy input was applied.

(2) At the proper E, the cross-sectional microstructure along thedepth of deposited layer demonstrated variant transforma-tions of grain morphology due to the change of temperaturegradient and solidification rate. At the same distance fromthe interface, the transformation of grain morphology delayedwith the applied E increasing, owing to the decrease of Gc in-duced by the accelerated diffusion coefficient within the mol-ten pool.

(3) Submicron TiO2 particles with unique flower-like structurewere obtained when the applied laser energy E ≤ 96 kJ/m.The formation mechanism of TiO2 submicronflower was pres-ent, mainly owing to the preferred growth of TiO2 along theb001N direction.

(4) As a proper laser energy E was set, the deposited layerachieved a relatively high microhardness with an average ofHV0.2 790 while the substrate material showed a low averagevalue of 280 HV0.2. Furthermore, the much smoother wornsurface with shallow grooves as well as a considerably lowerCOF of 0.4 was obtained, in comparison to the substrate mate-rial. The formation of Fe–Ti and Fe–Ni intermetallic as well asdispersive oxide particles were the main factors for the aboveimprovement.

Fig. 11. The coefficients of friction (COFs) as well as the partially corresponding worn surface morphologies of the substrate and the deposited layer obtained at different Es of 44 kJ/m,72 kJ/m, 96 kJ/m, 120 kJ/m and 144 kJ/m during the wear test.

52 C. Ma et al. / Surface & Coatings Technology 291 (2016) 43–53

Acknowledgments

The authors gratefully acknowledge the financial support fromthe National Natural Science Foundation of China (Nos. 51322509and 51575267), the Top-Notch Young Talents Program of China, theOutstanding Youth Foundation of Jiangsu Province of China (No.BK20130035), the Program for New Century Excellent Talents inUniversity (No. NCET-13-0854), the Science and Technology SupportProgram of Jiangsu Province (the industrial part), Jiangsu ProvincialDepartment of Science and Technology of China (No. BE2014009-2),the 333 Project (No. BRA2015368), Science and Technology Foundationfor Selected Overseas Chinese Scholar, Ministry of Human Resourcesand Social Security of China, the Aeronautical Science Foundation ofChina (No. 2015ZE52051), the Shanghai Aerospace Science and Tech-nology Innovation Fund (No. SAST2015053), the Fundamental ResearchFunds for the Central Universities (Nos. NE2013103 and NP2015206),and the Priority Academic Program Development of Jiangsu Higher Ed-ucation Institutions, the Foundation of the Graduate Innovation Centerin NUAA and the Fundamental Research Funds for the Central Universi-ties (No. kfjj20150605).

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