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Ti2Ni based laser cladding composite coatings
Optics and Laser Technology 97 (2017) 379–389
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Effect of heat treatment on residual stress and wear behaviors of the TiNi/Ti2Ni based laser cladding composite coatings Yang-Feng Tao, Jun Li ⇑, Ying-Hao Lv, Lie-Feng Hu School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China
a r t i c l e
i n f o
Article history: Received 13 January 2017 Received in revised form 26 April 2017 Accepted 18 July 2017
Keywords: Laser cladding Cracking susceptibility Residual stress Wear resistance Nanoindentation
a b s t r a c t The TiNi/Ti2Ni based composite coatings reinforced by TiC and TiB2 were prepared on Ti6Al4V at different circumstance temperatures (25 °C, 400 °C, 600 °C, and 800 °C) by laser cladding, then were preserved for 3 h. Macromorphologies and microstructures of the coatings were examined through an optical microscope (OM), an X-ray diffractometer (XRD), a scanning electron microscope (SEM), and an energy dispersive spectrometer (EDS). Residual stresses along the depth direction of the coatings were measured by the nanoindentation method, and wear behaviors of the coatings were also investigated using an ultrafunctional wear testing machine. Results showed that the coatings were mainly composed of TiNi/ Ti2Ni as the matrix and TiC/TiB2 as the reinforcement. A small amount of Cr2Ti was formed in the coatings prepared at 400 °C and 600 °C. Besides that, Ti3Al was also observed in the coating prepared at 800 °C. The tensile stress existed in the coatings prepared at 25 °C, 400 °C and 600 °C when the coating prepared at 800 °C was regarded as the stress-free reference. The average residual stress in the surface of coating prepared at 25 °C reached the largest value of about 2.79 GPa and presented a decreasing tendency with increasing the circumstance temperature (1.03 GPa at 400 °C, 0.52 GPa at 600 °C, and 0 GPa at 800 °C). It revealed that the rise in circumstance temperature contributed to the reduction in cracking susceptibility in the laser cladding coating. However, the wear volumes of the coatings were increased with increasing the circumstance temperature (0.1912 mm3 at 25 °C, 0.2828 mm3 at 400 °C, 0.3732 mm3 at 600 °C, and 0.6073 mm3 at 800 °C) due to the weakening in strain-hardening effect and the reduction in reinforcement density. The wear mechanism of the coatings was transformed from the single brittle-debonding into the combination of micro-cutting and brittle-debonding when the circumstance temperature was changed from room temperature to high temperature. The suitable circumstance temperature should be 600 °C, at which a comparatively high wear resistance was maintained on the premise that the residual stress was effectively relieved. Ó 2017 Published by Elsevier Ltd.
1. Introduction Titanium alloy is widely applied in aerospace, medical fields, marine and other industries due to its low density, high strength-to-weight ratio, excellent corrosion resistance, and the other superior performances. Especially in the aerospace field, the application of titanium alloy in aircraft components can endow the airplane with high maneuverability, high reliability, and long life. In the early 50 s, titanium alloy was used to fabricate some structural components subjected to low load in the military aircraft, such as thermal baffle, fairing, and airbrake [1,2]. At present, the application of titanium alloy is further expanded toward some structural components suffering force, such as flap, girder, bulk⇑ Corresponding author. E-mail addresses: [email protected] (Y.-F. Tao), [email protected] (J. Li), [email protected] (Y.-H. Lv), [email protected] (L.-F. Hu). http://dx.doi.org/10.1016/j.optlastec.2017.07.029 0030-3992/Ó 2017 Published by Elsevier Ltd.
head, and junction. However, further industrial applications of titanium alloy in some important structural components are limited owing to its low surface hardness and poor wear resistance. Therefore, surface modification of titanium alloy becomes a new research hotspot [3]. Surface properties of titanium alloy can be improved by surface modification technologies, such as plasma spraying [4], physical vapor deposition [5], chemical vapor deposition [6], nitriding and carburizing [7,8]. Among those methods, laser cladding is an effective and viable surface modification technique for improving wear resistance of titanium alloy by producing the ceramic-metal composite coating [9,10]. Sun et al. [11] fabricated a composite coating on Ti6Al4V by laser cladding the powder of NiCrBSi and TiC. The result showed that the wear weight loss of the coating was 11.4% that of Ti6Al4V in atmosphere and 47.9% in vacuum. Meng et al. [12] prepared two coatings on Ti6Al4V by laser cladding. The NiCrBSi + 5 wt.% B4C powder and NiCoCrAlY powder were selected as the cladding materials, respectively. Wear
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resistance of the substrate was improved significantly since the average hardness of the coatings was 2–4 times that of Ti6Al4V. Savalani et al. [13] produced some TiC reinforced Ti matrix composite coatings on pure titanium by laser cladding the mixed powders of Ti and carbon-nanotube with different contents. They found that the coatings with higher carbon-nanotube presented better wear resistance. Li et al. [14] synthesized a TiN and TiB reinforced metal matrix composite coating with a Ti/h-BN powder mixture as the cladding material by laser cladding. The result showed that the composite coating with a hardness of 800 HV-1200 HV presented better wear resistance. Although the laser cladding technique can obviously improve surface properties of titanium alloy such as hardness and wear resistance, the coating remains a serious cracking susceptibility due to the existence of high residual stress, which vastly limits the industrial application of laser cladding. The residual stress in the laser cladding coating can be divided into thermal stress, structural stress and restrain stress. Among them, the thermal stress is predominant, which mainly depends on the difference in temperature between environment and cladding. The uneven distribution temperature results in a high thermal gradient in the coating. The uneven expansion and contraction occur along with the laser cladding process. Thus, the high residual stress is produced during laser cladding [15]. Many methods have been used to reduce the cracking susceptibility of the laser cladding coating such as the optimization of processing parameters [16], preheating and postheat treatments [17,18], and adding the alloying powders [19]. Among those methods, preheating and postheat treatments can effectively reduce the temperature gradient in the coating (especially between coating and substrate). The thermal stress can be decreased significantly. Therefore, this method is widely used to reduce the cracking susceptibility of the laser cladding coating. Vahid et al. [20] found that the crack formation in laser cladding the hardfacing alloy Stellite 1 on AISI-SAE 4340 steel was prevented through locally preheating the substrate prior to the deposition process. The result showed that the preheated sample revealed more uniform structure and was less prone to cracking. Fallahi et al. [21] introduced an efficient approach to reduce the residual stress in the welding process by a preheating treatment. The residual stress was decreased when the preheating temperature was increased from 25 °C to 400 °C. When the coating is prepared at a high circumstance temperature, the residual stress in the coating is released and the phase transformation possibly happens. These changes will result in the elimination in strain-hardening effect and the evolution in coating microstructure. As a result, wear resistance of the coating will be inevitably influenced at the same time. Therefore, it is necessary to study the influence of preheating and postheat treatments on wear resistance of the coating with a low residual stress. The investigation will contribute to confirming the appropriate treatment temperature, in which not only the cracking susceptibility is reduced, but also excellent wear resistance is maintained. However, there are hardly investigations about the issue. In this study, the Ti2Ni/TiNi based coatings were prepared on Ti6Al4V at different circumstance temperatures (25 °C, 400 °C, 600 °C and 800 °C) by laser cladding, then were preserved for 3 h. The relationship between cracking susceptibility and wear resistance of the coatings was investigated in detail. Finally, an appropriate temperature was determined to obtain the best comprehensive mechanical properties of the coatings.
2. Experimental procedures Ti6Al4V [composition (in wt.%): 6.5 Al, 4.26 V, 0.1 C, balance Ti] which was cut into the samples with a dimension of 30 mm ⁄ 20
mm ⁄ 10 mm was used as the substrate. F102 Ni-based alloy powder [composition (in wt.%): 75 Ni, 1 C, 16 Cr, 3.5 B, 4.5 Si] was selected as the cladding material. The substrate surface was brushed with a binder (4% polyvinyl alcohol) and then placed in a cubic model with a dimension of 30.2 mm ⁄ 20.2 mm ⁄ 10.8 mm. A space with a precise height of 0.8 mm was formed above the substrate. The powder was filled into this the space and compacted to form a preplaced layer at 30 MPa for 3 mins using a tablet machine. This method not only ensures the uniformity in thickness of the preplaced layer, but also reduces its porosity. Before laser cladding, the samples were preheated to 400 °C, 600 °C and 800 °C, and were maintained for half an hour in a resistance furnace. After that, the furnace lid was opened and laser cladding was directly carried out on the samples placed in the furnace. The circumstance temperature during laser cladding was nearly maintained at the initially set temperature (400 °C, 600 °C, and 800 °C), because the laser cladding time was very short (about 15 s) and the furnace was still in work during laser cladding. Laser cladding was performed using an YLS-5000 fiber laser processing system with an applied power of 3 KW, a spot diameter of 6 mm, and a scanning speed of 10 mm/s. The parameter was optimized by ensuring a good interface fusion and a smooth coating surface. The sample without preheating (25 °C) as the contrast was prepared at the same processing parameter, and the obtained coating was named as coating I. The coatings prepared at 400 °C, 600 °C and 800 °C were named as coatings II, III and IV. After laser cladding, the samples were preserved at 400 °C, 600 °C and 800 °C for 3 hours, respectively. Phase constituents of the coatings were analyzed through a PANalytical X’ Pert Pro X-ray diffractometer (XRD). Macromorphologies of the coatings were observed with a VHX-600 K optical microscope (OM). Microstructural characterization was conducted on the cross-sections of coatings using an S-3400 scanning electron microscope (SEM) coupled with a GENESIS EDAX energy dispersive spectrometer (EDS). Microhardness of the cross-sections of coatings was measured by a HXD-1000TMSC/LCD microhardness tester with a load of 200 gf applied for 15 s. Nanoindentation tests were performed using a TriboIndenterÒ (Hysitron Corporation, USA) with a diamond Berkovich indenter. The indenter was also used as an AFM tip to imaging the indented surface after the tests. A peak load of 10 mN was selected in the tests, and the loading and unloading times were set at 5 seconds each and the maximum load was also maintained for 5 seconds. The first nanoindentation was obtained at the zone of 0.1 mm distance from the coating surface. The next indentations were prepared at different zones with the same distance interval (0.4 mm). Three zones in every coating were selected for the tests. Room-temperature dry friction tests were performed to evaluate wear resistance of the coatings using a CFT-I ultra-functional wear testing machine in the ball-on-disc reciprocating motion mode. The sliding time was 180 min, the applied load was 30 N, and the reciprocating sliding speed was 5 m/min. YG6 balls with a diameter of 6 mm were used as the counterparts. Wear surface morphologies of the coatings were observed with SEM. Wear volumes of the coatings were measured using a wear profiler.
3. Results and discussion 3.1. Macromorphologies Fig. 1 shows the macromorphologies of the coatings prepared at different circumstance temperatures. The profiles of the coatings are very similar, presenting the convex shape. No obvious cracks and unmelted particles are observed. A good metallurgical bonding is formed between coatings and substrate. The width (6.21 mm at
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Fig. 1. Morphologies of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
25 °C, 7.00 mm at 400 °C, 7.08 mm at 600 °C and 7.20 mm at 800 °C) and thickness (1.98 mm at 25 °C, 2.10 mm at 400 °C, 2.44 mm at 600 °C and 2.84 mm at 800 °C) of the four coatings present an increase tendency with the increase in circumstance temperature. Dilution rate g can be calculated as follows [22]:
g¼
S1 100% S1 þ S2
ð1Þ
where S1 is the molten area of the substrate, and S2 is the area of the coating above the surface of the substrate. The dilution rate is determined to be 64.40% (25 °C), 70.21% (400 °C), 72.53% (600 °C) and 74.29% (800 °C). The dilution rate of the coatings is increased with the increase in circumstance temperature due to more extra energy provided in the cladding system. 3.2. Microstructural characterization Fig. 2 illustrates the XRD spectra of the coatings prepared at different circumstance temperatures. The spectrum of the coating prepared at 25 °C is very simple. Main diffraction peaks are in accordance with the Joint Committee on Powder Diffraction Standards (JCPDS) cards (No. 03-065-5537 for TiNi, No. 00-018-0898 for Ti2Ni, No. 03-065-8698 for TiB2, and No. 03-065-8804 for TiC). The three strongest lines of every phase are in good agreement with those in the JCPDS cards. Therefore, the coating is composed of TiNi, Ti2Ni, TiB2, and TiC. When the preheating temperature is increased to 400 °C, two extra diffraction peaks at 2h = 39.1° and 2h = 39.6° are observed and the intensity of a diffraction peak at 2h = 42.7° is significantly increased. It can be confirmed that a new phase of Cr2Ti was in situ synthesized in the coating prepared at 400 °C. The XRD spectrum of the coating prepared at 600 °C is highly similar to that of the coating prepared at 400 °C, indicating that their phase constituents are same. As the preheating temperature was further increased to 800 °C, a significant change is visible. Three new diffraction peaks appear at 2h = 47.8°, 2h = 71.0° and 2h = 76.7°. In terms of the JCPDS cards, the other new phase of Ti3Al was formed in the coating prepared
at 800 °C. Other than that, the intensities of two diffraction peaks (2h = 41.5° and 2h = 42.7°) also present the increasing tendency when compared with those in Fig. 2b and c. Considering that the contents of Ni and C are same in the four coatings, the increase in intensity of the two diffraction peaks should be attributed to the increase in volume fraction of Ti2Ni and Cr2Ti. As mentioned above, the dilution rate of the coatings is increased with the increase in circumstance temperature (64.40% at 25 °C, 70.21% at 400 °C, 72.53% at 600 °C, and 74.29% at 800 °C). The increase in dilution rate means that more Ti enters into the molten pool during laser cladding. As a result, some compounds rich in Ti (Cr2Ti and Ti3Al) are formed successively with the increasing in circumstance temperature, and the volume fractions of three compounds rich in Ti (Ti2Ni, Cr2Ti and Ti3Al) also show the increasing tendency. Fig. 3 shows the Backscattered Electron (BSE) images of representative microstructures of the four coatings. A large number of acicula-shaped structures are distributed uniformly in the matrix of coating I. As indicated in the XRD spectra, the reinforcements in the four coatings mainly include TiB2 and TiC. When the circumstance temperature is increased to 400 °C and 600 °C, the microstructure of coatings II, III is similar to that in coating I. However, the density of acicula-shaped structures exhibits the downward trend. With the increase in circumstance temperature (800 °C), the density of the reinforcements is further decreased. Moreover, the morphology of the reinforcements is transformed from coarse acicula shape into fine feather shape. It is well known that the growth of the reinforcements strongly depends on the atomic content in the molten pool. The dilution rate of coating III is highest among the four coatings, which indicates that the concentration of B and C is also lowest in the molten pool. During solidification of the molten pool, the growth of the reinforcements in coating III will be restricted greatly when compared with the other three coatings. As a result, the reinforcements in the coating prepared at 800 °C are finer. The microstructure in a high-magnification image further reveals that coating I is composed of six phases (shown in Fig. 4), corresponding to cupped light gray phase (phase 1), protuberant gray white phase (phase 2), grey strip-shaped/spherical particles (phase 3), block acicula-shaped particles (phase 4), rodlike parti-
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Fig. 2. X-ray diffraction spectra of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
Fig. 3. Microstructures of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
cles (phase 5), and equiaxed particles (phase 6). Phases 1 and 2 with a high volume fraction can be regarded as the matrix, phases 3, 4, 5 and 6 should be the in situ synthesized reinforcements. The rodlike (phase 5) and equiaxed particles (phase 6) as the primary phase exist independently in the matrix. However, the comparatively fine grey strip-shaped/spherical particles (phase 3) are uniformly distributed in the block acicula-shaped particles (phase 4), which can be identified as a eutectic structure. The compositions of those phases with different morphologies were analyzed by EDS, and the results are shown in Table 1. Phases 1 and 2 are rich in titanium and nickel, accompanied with a small amount of Cr, Al and Si, which should be two kinds of Ti-Ni compounds. Ti
and Ni atoms are easily replaced by (Cr and Al) and Si atoms, respectively due to their similar atomic radii and electronegativities. The atomic percentage of Ti, Cr and Al in phase 1 is 48.00 at. %, and that of Ni and Si is 49.70 at.%. Thus the cupped light gray phase (phase 1) can be identified as the TiNi solid solution. Similarly, the protuberant gray white phase (phase 2) can be regarded as Ti2Ni since the atomic ratio of (Ti + Al + Cr) and (Ni + Si) is about 1.54:1. The grey strip-shaped/spherical particles (phase 3), rodlike particles (phase 5) and equiaxed particles (phase 6) with very similar chemical compositions are rich in titanium and carbon. The block acicula-shaped particles (phase 4) mainly consists of titanium (15.37 at.%) and boron (84.08 at.%). Combined with the
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Fig. 4. Typical microstructure of coating I.
Table 1 EDS analyses (in at.%). Coatings
Phase
Ti
C
B
Ni
Cr
Al
Si
Phase information
I
1 2 3 4 5 6
32.49 39.06 48.34 15.37 40.68 51.65
1.24 1.54 51.66 0.55 59.32 48.35
– – – 84.08 – –
48.12 26.24 – – – –
4.24 13.66 – – – –
11.27 0.71 – – – –
1.58 8.46 – – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
II
1 2 3 4 5 6
41.01 34.74 61.72 13.56 41.52 50.63
2.16 2.18 33.54 – 58.48 49.37
– – – 86.44 – –
42.84 27.91 2.36 – – –
4.80 17.49 0.46 – – –
6.58 5.04 0.52 – – –
1.26 8.85 0.84 – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
III
1 2 3 4 5 6
40.46 36.23 60.30 14.05 45.92 48.69
1.13 2.14 39.70 – 54.08 51.31
– – – 85.95 – –
45.29 25.50 – – – –
4.10 16.04 – – – –
6.85 5.45 – – – –
1.00 10.47 – – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
IV
1 2 3 4 5 6
51.07 41.93 51.58 22.41 44.35 47.33
2.81 3.18 48.42 – 55.65 52.67
– – – 77.59 – –
26.86 26.11 – – – –
4.80 10.23 – – – –
8.73 9.70 – – – –
4.29 5.86 – – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
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XRD analyses, phases 3, 5, 6 can be identified as TiC and phase 4 should be TiB2. The compositions of six phases in the other coatings are very similar to those in coating I (shown in Table 1). The reinforcements in the composite coatings are TiC and TiB2. The microstructural evolution of the coating can be interpreted through a TiB2-TiC phase diagram [23]. The eutectic structure is formed when the mole fraction of TiC is 0.2175. The mole fraction of TiC in the molten pool can be calculated as about 0.3432 in this study, which indicates that the content of TiC is located in the hypereutectic area. Therefore, TiC as the primary phase is first precipitated from the liquid and then grows into the rodlike or equiaxed particles. TiC has a B1 crystal structure and the geometry of TiC crystal cell presents the octahedral structure. The TiC crystal is inclined to grow along the h100i directions preferentially due to its face centered cubic structure [24]. When the degree of supercooling in the molten pool is small, the octahedral TiC crystal uniformly branches in the six h100i directions with the similarly same growth speed, resulting in the formation of equiaxed TiC particles. However, the crystal in a specific direction may grow faster than the other directions during subsequent solidification because of the differences in temperature and concentration distributions in different directions [25]. As a result, the rodlike TiC particles are formed. With the decrease in temperature in the molten pool, a eutectic structure composed of TiB2 and TiC is formed by heterogeneous nucleation and grows into the block acicula-shaped structure [26]. Afterwards, Ti2Ni/TiNi are formed.
3.3. Nanoindentation tests Fig. 5 shows the load-depth curves of the coatings prepared at different circumstance temperatures. According some data obtained from Fig. 5, some mechanical properties (such as the residual stress) of the coatings can be extracted by the OliverPharr method. The key parameter in the calculation is the contact area, which is directly influenced by the ratio of the final indentation depth (hf) and the maximum indentation depth (hmax). The ratio reflects the degree of pile-up deformation in the coating. A serious pile-up deformation will generate around the indentation edge when hf/hmax is close to 1. As a result, the area deduced from the load-depth curves is far less than the true area by as much as 60% [27]. This leads to a serious overestimation of mechanical properties of the coating, even a significant error (exceeding fifty percent) may be produced. When hf/hmax < 0.7, very little pile-up occurs and a comparatively accurate result can be obtained. In this study, the values of hf/hmax at different zones of the coatings are listed in Table 2. It is clear that all values are lower than 0.7. Fig. 6 shows the two-dimensional and three-dimensional morphologies of a typical nanoindentation of coating III. No obvious pile-up deformation is observed in the zone around the nanoindentation. The Suresh model is commonly used to distinguish and calculate the residual stress of the material, which is mainly based on the difference in maximum depth between a stress-free material
Fig. 5. Load-depth curves of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
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Table 3 Residual stress at different zones in the coatings (GPa).
Coatings
0.1 (mm)
0.5 (mm)
0.9 (mm)
Reference
Coatings
0.1 (mm)
0.5 (mm)
0.9 (mm)
I II III IV
0.61 0.58 0.59 0.56
0.50 0.62 0.60 0.58
0.57 0.57 0.59 0.58
Coating IV
I II III IV
2.51 1.72 1.03 0.00
2.56 1.94 1.05 0.00
3.27 0.58 0.52 0.00
Coating III
I II III IV
1.78 0.81 0.00 1.24
1.60 1.10 0.00 1.31
1.96 0.04 0.00 0.49
Coating II
I II III IV
1.13 0.00 0.92 2.34
0.32 0.00 1.39 3.05
2.00 0.00 0.04 0.53
Coating I
I II III IV
0.00 1.41 2.51 4.26
0.00 0.28 1.74 3.48
0.00 2.44 2.30 1.67
and a material with residual stress. Therefore, the selection of a stress-free material is a very key factor for the calculation of the residual stress. Considering that the microstructure of the laser cladding coating is very complex, it is difficult to prepare a stress-free coating with similar microstructure by the other methods. A feasible method is to select a material or a specific zone in a material with a comparatively low stress as the stress-free reference among all materials, which had been used widely in previous studies. Zhu et al. [27] studied the residual stress in the plasmasprayed FeCrBSi coating by the nanoindentation method. As the residual stress in the interface between FeCrBSi coating and Ni/Al transition layer was close to zero, the interface was regarded as the stress-free reference. Zhang et al. [28] investigated the residual stress in 304 stainless steel by the nanoindentation method. The annealed stainless steel was selected as the stress-free sample to measure the residual stress in the unannealed sample. The residual stress was about 381 MPa, which was very close to the result tested by the XRD method (350 ± 23 MPa). Wang et al. [29] measured the residual stress in the plasma sprayed Fe-based coating using the nanoindentation method. A coating with the 0.5 mm thick surface removed by the wire-electrode cutting method was selected as the stress-free coating since the residual stress was released fully. According to the fundamental theory of materials science, the residual stress in the coating can be reduced by reducing the temperature difference between environment and cladding. Therefore, the coating prepared at 800 °C was selected as the stress-free coating in our study, In order to verify the rationality, the residual stress in the four coatings was calculated when coatings I, II, III and IV were selected as the stress-free coating, respectively. As shown in Table 3, the residual compressive stress exists in some coatings when coatings I, II and III are selected as the stress-free coating. The residual tensile stress exists in the three coatings when coating IV is regarded as the stress-free coating. Some studies had confirmed that the residual tensile stress was produced easily in the laser cladding coating due to the rapid
heating and cooling process during laser cladding [29,30]. Therefore, it is appropriate to select coating IV as the stress-free coating. The residual stresses in the other coatings were calculated by the following equations [31]: For the tensile residual stress:
r¼H 1
2
h0 h
! ð2Þ
2
For the compressive residual stress:
r¼
2
H h0 1 sin a h2
! ð3Þ
where h and h0 are the penetration depth in the coatings with and without the residual stress, respectively, H is the hardness of the coatings, and a is related to the indentation angle of the indenter (a = 24.7 for a Berkovich indenter). Fig. 7 indicates the load-depth curves obtained at different zones of the four coatings. The maximum depth in coating I always reaches the maximum value. At the zones with 0.1 mm and 0.5 mm distance from the surfaces of the coatings, the maximum depth presents the decreasing trend with the increase in circumstance temperature (shown in Fig. 7a and b). At the zones with
Fig. 6. Two-dimensional and three-dimensional morphologies of a typical nanoindentation of coating III.
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Fig. 7. Load-depth curves for the coatings: (a) 0.1 mm, (b) 0.5 mm and (c) 0.9 mm.
about 0.9 mm distance from the surfaces of coatings I, II and III (shown in Fig. 7c), the maximum depth is approximately same. When the same load is applied to the coatings, the different indentation depth is produced, which should be attributed to the existence of the residual stress in the coatings. The tensile stress tends to facilitate the indentation process. However, the compressive stress will produce the contrary effect in the indentation formation. Therefore, it is concluded that the tensile stress exists in coatings I, II, and III when coating IV is regarded as the stressfree reference. The residual tensile stress in the coatings was calculated by the Suresh model (shown in Table 3). The residual tensile stress at the zones with different distance from the surface of coating I always reaches the maximum value (2.51 GPa in 0.1 mm, 2.56 GPa in 0.5 mm and 3.27 GPa in 0.9 mm). At the zones with 0.1 mm distance from the surfaces of the coatings, the residual tensile stress is reduced about 25.1% (1.72 GPa) and 40.3% (1.03 GPa) with increasing the circumstance temperature to 400 °C and 600 °C. At the zones with 0.5 mm distance from the surfaces of the coatings, the residual tensile stress in the coatings is reduced about 24.4% (1.94 GPa for 400 °C) and 45.6% (1.05 GPa for 600 °C). However, the residual tensile stress in coatings II and III approaches zero at the zones with 0.9 mm distance from the surfaces of the coatings. It can be concluded that the residual stress in the coatings prepared at high circumstance temperatures could be reduced obviously in whole. Fig. 8 shows the microhardness distribution throughout the coatings. The microhardness shows a downward trend along the
depth direction of the cross-sections of the coatings. Considering the change in microhardness, three regions can be clearly identified, corresponding to the coating, the transition zone and the substrate. The average microhardness values of coating I at the three zones are 1062.2 HV0.2, 624.2 HV0.2, and 350.8 HV0.2. The average microhardness value of the coating is about 3 times that of the sub-
Fig. 8. Microhardness distribution across the cross sections of the coatings.
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strate due to the uniform distribution of a large number of highhardness TiC (30 GPa) [32] and TiB2 (34 GPa) [33] particles within TiNi and Ti2Ni intermetallic compounds. The existence of the transition zone contributes to forming a good metallurgical bonding between coating and substrate. Moreover, the transition zone plays a buffer effect between hard coating and soft substrate. The coating and the substrate have different elasticity modulus and thermal expansion coefficient, which leads to different deformation when the coating is subject to the external force or temperature variation in service. As a result, the huge stress concentration will generate in the junction interface between coating and substrate. The crack will initiate and propagate when the stress exceeds the bonding strength. Therefore, the transition zone can effectively reduce the cracking sensibility of the coating. The average microhardness values of coatings II, III and IV are 1031.2 HV0.2, 1028.2 HV0.2, and 954.7 HV0.2, respectively. It is clear that the average microhardness values of the coatings are reduced as increasing the circumstance temperature. The decline in microhardness and residual stress of the coatings has the close relationship with the changes in microstructure (especially atomic migration). The high residual stress usually exists in the coating fabricated by laser cladding due to its rapid heating and cooling characteristics. The strainhardening effect resulting from the residual stress will further improve the microhardness of the coating. The number of defects in the coating is reduced due to the rapid atomic migration when the coating is prepared and maintained at high circumstance temperature. Accompanied with this, the residual stress is released and the microhardness is decreased accordingly. On the other hand, the reduction in reinforcement density can also cause the decrease in microhardness to a certain extent. 3.4. Wear resistance Wear profiles of the coatings are shown in Fig. 9. The wear volume is 0.1912 mm3 for coating I. The wear volumes are increased by nearly 47.9% (0.2828 mm3 for coating II), 95.2% (0.3732 mm3 for coating III) and 220% (0.6073 mm3 for coating IV) with the increase in circumstance temperature. Fig. 10 shows the surface morphologies of coatings I, II, III and IV after the wear tests. Many spalling holes are observed on the worn surface of coating I (shown in Fig. 10a). The other zones except the spalling holes are comparatively smooth. Coatings II, III and IV present the similar worn surface morphology (shown in Fig. 10b–d). However, the number of the spalling hole on the worn surfaces of the coatings is decreased with increasing the circumstance tem-
Fig. 9. Wear profiles of the coatings.
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perature. The distribution and morphological characteristic of the hole resulting from the spalling of wear debris can be used to reveal the formation mechanism of the debris. The spalling hole generally generates in the center or edge of the reinforcements and hardly generates in the matrix. During sliding friction, the difference in elasticity modulus between reinforcement and matrix leads to their different deformation when a load is applied to the contact surface. The stress concentration will generate in the interface. The crack may initiate and propagate along the interface, finally into the matrix when the stress exceeds the interface bonding strength. Moreover, the crack may generate within the reinforcement and extend toward the matrix due to its low toughness. The bottom of the spalling hole is actually the fracture surface, which records the details of crack propagation. The riverlike pattern can be clearly observed at the bottom of the spalling hole, which presents the typical brittle-fracture characteristic. The images with the high magnification clearly indicate that the morphology of the fracture surface of coating I well accords with that of the typical intergranular fracture (shown in Fig. 11a). Many grains can be clearly observed, and the whole surface is comparatively smooth due to the insignificant change in contrast. A small number of white tearing edges can be only visible in the grain boundaries. It can be concluded that the plastic deformation does not almost occur during crack propagation. That is to say, the crack propagates with less resistance. The large residual tensile stress in coating I will further promote the crack initiation and propagation. The impurity elements (such as O, N) have no enough time to escape from the molten pool during laser cladding due to its rapid heating and cooling characteristic. They may aggregate in the grain boundaries of TiNi and Ti2Ni, which reduce the bonding strength of the boundaries. The crack tends to initiate and propagate along the grain boundaries during wear, resulting in the intergranular fracture. Therefore, wear of coating I is mainly dominated by brittledebonding. As shown in Fig. 11b, the surface of the spalling hole in coating III is rougher when compared with that of coating I, which indicates that the comparatively large plastic deformation occurs during crack propagation. The plastic deformation will consume the extra energy and hinder the crack propagation. On the other hand, the decrease in residual tensile stress will further restrain the crack propagation. Therefore, the brittle-debonding is relieved with increasing the circumstance temperature. Moreover, some slight furrows and scratches are clearly observed on the worn surface of coating III (as shown in Fig. 11b). It can be concluded that wear mechanism of the coatings is transformed from the single brittle-debonding into the combination of the micro-cutting and the brittle-debonding when the processing temperature is changed from room temperature to high temperature. The phenomenon is closely related to the changes in hardness and residual stress of the coatings. As analyzed above, the debris is mainly formed by two mechanisms corresponding to micro-cutting and brittle-debonding. The resistance to micro-cutting mainly depends on hardness of the coating, the resistance to brittle-debonding is closely related to plasticity and residual stress of the coating. With the change in preheating temperature, the resistance to micro-cutting and brittle-debonding of the coating will present the regular change. Among the four coatings, the hardness of coating I is comparatively high due to the high density of the reinforcements and the strong strain-hardening effect, indicating that the coating has the powerful resistance to micro-cutting. The coating with high hardness usually possesses the comparatively low plasticity. Moreover, a very high residual tensile stress exists in coating I without preheating. The mutual effect of the two factors will significantly accelerate the crack initiation and propagation in coating I when it is in service. That is to say, coating I exhibits very high cracking susceptibility and very less resistance to brittle-debonding. When the
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Fig. 10. Morphologies of the wear surfaces of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
Fig. 11. Typical morphologies of the wear surface of the coatings: (a) 25 °C and (b) 600 °C.
preheating temperature is increased to 400 °C and 600 °C, the hardness of the coating is reduced from HV 10620.2 to about 1030 HV0.2 due to the reduction in density of reinforcements and the weakening in strain-hardening. It indicates that the two coatings present the weaker resistance to micro-cutting than coating I. However, their resistance to brittle-debonding is improved to a certain extent along with the increase in plasticity and the decrease in residual stress. The combination of the two effects (micro-cutting and brittle-debonding) results in about 0.3280 mm3 loss of coatings II and III, which is higher than that of coating I. Similarly, when the preheating temperature is increased to 800 °C, the resistance to micro-cutting of coating IV is further reduced, accompanied with the increase in resistance to brittle-debonding.
The wear volume of coating IV is further enhanced to 0.6073 mm3. According the change in wear volume with the increase in preheating temperature, it can be concluded that the microcutting plays the important role in debris formation. The debris is immediately produced with the movement of the counterpart during the micro-cutting process, while the brittle-debonding process involves crack initiation, propagation and debonding, which needs more time. Thus, the micro-cutting is easier to produce debris, when compared with the brittle-debonding. Based the above analyses, it can be concluded that laser cladding at the high temperature is useful to reduce cracking susceptibility in the coatings. However, wear resistance of the coatings will be decreased to a certain extent. When the preheating temperature
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exceeds 600 °C, the wear volume of the coating (800 °C) is increased about 220% when compared with that without preheating. Wear resistance of the coating is weakened significantly, which greatly reduces its service life. When the preheating temperature is less than 600 °C, the wear volume of the coating (400 °C) is only increased about 47.9% when compared with that without preheating. The coating still presents the excellent wear resistance. However, it also has a high cracking susceptibility since a comparatively high residual tensile stress exists. Comparatively speaking, the wear volume of the coating prepared at 600 °C is slightly lower than that of the coating prepared at 400 °C. However, the residual stress is further reduced about 40%. Therefore, the suitable preheating temperature should be 600 °C, at which the comparatively high wear resistance is maintained on the premise that the residual stress is relieved. 4. Conclusions (1) TiNi/Ti2Ni-based composite reinforced by TiC and TiB2 were prepared on the Ti6Al4V substrate at different circumstance temperatures (25 °C, 400 °C, 600 °C and 800 °C) by laser cladding, then were preserved for 3 h. (2) The tensile stress existed in the coatings when the coating prepared at 800 °C was regarded as the stress-free reference. The average residual stress in the surface of coating prepared at 25 °C reached the largest value of about 2.79 GPa and presented the decreasing tendency with increasing the circumstance temperature (1.03 GPa at 400 °C, 0.52 GPa at 600 °C and 0 GPa at 800 °C). (3) The wear volumes of the coatings prepared at 25 °C, 400 °C, 600 °C and 800 °C were increased with the increase in circumstance temperature (0.1912 mm3 at 25 °C, 0.2828 mm3 at 400 °C, 0.3732 mm3 at 600 °C and 0.6073 mm3 at 800 °C). The brittle-debonding was predominant in the debris formation of the coating prepared at 25 °C. Wear mechanism was transformed into the combination of the microcutting and brittle-debonding with the increase in processing temperature. The micro-cutting caused a more serious wear loss of the coating, when compared with the brittledebonding. (4) The processing temperature should be 600 °C, at which not only the high wear resistance was maintained, but also the cracking susceptibility was reduced.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (51471105), ‘‘Shu Guang” project of Shanghai Municipal Education Commission and Shanghai Education Development Foundation (12SG44) and ‘‘Graduate innovation” Project of Shanghai University of Engineering Science (15KY0504). References [1] M’S. Rachid, D.s. Axinte, S.L. Soo, C. Nobel, H. Attia, G. Kappmeyer, S.F. Engin, W. M. Sim, High performance cutting of advanced aerospace alloys and composite materials, Cirp. Ann.-Manuf. Techn. 64 (2015) 557–580. [2] Y. Yang, C.H Zhang, Y.J Dai, J.B Luo, Tribological properties of titanium alloys under lubrication of SEE oil and aqueous solutions, Tribol Int. 109 (2017) 40– 47. [3] S.K. Jayme, A.P. Todd, Development of strength-hardness relationships in additively manufactured titanium alloys, Mater. Sci. Eng., A. 693 (2017) 214– 224. [4] A.D. Jamaloeia, H.R. Salimijazia, H. Edrisa, J. Mostaghimib, Study of TLP bonding of Ti-6Al-4V alloy produced by vacuum plasma spray forming and forging, Mater. Des. 121 (2017) 355–366.
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Full length article
Effect of heat treatment on residual stress and wear behaviors of the TiNi/Ti2Ni based laser cladding composite coatings Yang-Feng Tao, Jun Li ⇑, Ying-Hao Lv, Lie-Feng Hu School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China
a r t i c l e
i n f o
Article history: Received 13 January 2017 Received in revised form 26 April 2017 Accepted 18 July 2017
Keywords: Laser cladding Cracking susceptibility Residual stress Wear resistance Nanoindentation
a b s t r a c t The TiNi/Ti2Ni based composite coatings reinforced by TiC and TiB2 were prepared on Ti6Al4V at different circumstance temperatures (25 °C, 400 °C, 600 °C, and 800 °C) by laser cladding, then were preserved for 3 h. Macromorphologies and microstructures of the coatings were examined through an optical microscope (OM), an X-ray diffractometer (XRD), a scanning electron microscope (SEM), and an energy dispersive spectrometer (EDS). Residual stresses along the depth direction of the coatings were measured by the nanoindentation method, and wear behaviors of the coatings were also investigated using an ultrafunctional wear testing machine. Results showed that the coatings were mainly composed of TiNi/ Ti2Ni as the matrix and TiC/TiB2 as the reinforcement. A small amount of Cr2Ti was formed in the coatings prepared at 400 °C and 600 °C. Besides that, Ti3Al was also observed in the coating prepared at 800 °C. The tensile stress existed in the coatings prepared at 25 °C, 400 °C and 600 °C when the coating prepared at 800 °C was regarded as the stress-free reference. The average residual stress in the surface of coating prepared at 25 °C reached the largest value of about 2.79 GPa and presented a decreasing tendency with increasing the circumstance temperature (1.03 GPa at 400 °C, 0.52 GPa at 600 °C, and 0 GPa at 800 °C). It revealed that the rise in circumstance temperature contributed to the reduction in cracking susceptibility in the laser cladding coating. However, the wear volumes of the coatings were increased with increasing the circumstance temperature (0.1912 mm3 at 25 °C, 0.2828 mm3 at 400 °C, 0.3732 mm3 at 600 °C, and 0.6073 mm3 at 800 °C) due to the weakening in strain-hardening effect and the reduction in reinforcement density. The wear mechanism of the coatings was transformed from the single brittle-debonding into the combination of micro-cutting and brittle-debonding when the circumstance temperature was changed from room temperature to high temperature. The suitable circumstance temperature should be 600 °C, at which a comparatively high wear resistance was maintained on the premise that the residual stress was effectively relieved. Ó 2017 Published by Elsevier Ltd.
1. Introduction Titanium alloy is widely applied in aerospace, medical fields, marine and other industries due to its low density, high strength-to-weight ratio, excellent corrosion resistance, and the other superior performances. Especially in the aerospace field, the application of titanium alloy in aircraft components can endow the airplane with high maneuverability, high reliability, and long life. In the early 50 s, titanium alloy was used to fabricate some structural components subjected to low load in the military aircraft, such as thermal baffle, fairing, and airbrake [1,2]. At present, the application of titanium alloy is further expanded toward some structural components suffering force, such as flap, girder, bulk⇑ Corresponding author. E-mail addresses: [email protected] (Y.-F. Tao), [email protected] (J. Li), [email protected] (Y.-H. Lv), [email protected] (L.-F. Hu). http://dx.doi.org/10.1016/j.optlastec.2017.07.029 0030-3992/Ó 2017 Published by Elsevier Ltd.
head, and junction. However, further industrial applications of titanium alloy in some important structural components are limited owing to its low surface hardness and poor wear resistance. Therefore, surface modification of titanium alloy becomes a new research hotspot [3]. Surface properties of titanium alloy can be improved by surface modification technologies, such as plasma spraying [4], physical vapor deposition [5], chemical vapor deposition [6], nitriding and carburizing [7,8]. Among those methods, laser cladding is an effective and viable surface modification technique for improving wear resistance of titanium alloy by producing the ceramic-metal composite coating [9,10]. Sun et al. [11] fabricated a composite coating on Ti6Al4V by laser cladding the powder of NiCrBSi and TiC. The result showed that the wear weight loss of the coating was 11.4% that of Ti6Al4V in atmosphere and 47.9% in vacuum. Meng et al. [12] prepared two coatings on Ti6Al4V by laser cladding. The NiCrBSi + 5 wt.% B4C powder and NiCoCrAlY powder were selected as the cladding materials, respectively. Wear
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resistance of the substrate was improved significantly since the average hardness of the coatings was 2–4 times that of Ti6Al4V. Savalani et al. [13] produced some TiC reinforced Ti matrix composite coatings on pure titanium by laser cladding the mixed powders of Ti and carbon-nanotube with different contents. They found that the coatings with higher carbon-nanotube presented better wear resistance. Li et al. [14] synthesized a TiN and TiB reinforced metal matrix composite coating with a Ti/h-BN powder mixture as the cladding material by laser cladding. The result showed that the composite coating with a hardness of 800 HV-1200 HV presented better wear resistance. Although the laser cladding technique can obviously improve surface properties of titanium alloy such as hardness and wear resistance, the coating remains a serious cracking susceptibility due to the existence of high residual stress, which vastly limits the industrial application of laser cladding. The residual stress in the laser cladding coating can be divided into thermal stress, structural stress and restrain stress. Among them, the thermal stress is predominant, which mainly depends on the difference in temperature between environment and cladding. The uneven distribution temperature results in a high thermal gradient in the coating. The uneven expansion and contraction occur along with the laser cladding process. Thus, the high residual stress is produced during laser cladding [15]. Many methods have been used to reduce the cracking susceptibility of the laser cladding coating such as the optimization of processing parameters [16], preheating and postheat treatments [17,18], and adding the alloying powders [19]. Among those methods, preheating and postheat treatments can effectively reduce the temperature gradient in the coating (especially between coating and substrate). The thermal stress can be decreased significantly. Therefore, this method is widely used to reduce the cracking susceptibility of the laser cladding coating. Vahid et al. [20] found that the crack formation in laser cladding the hardfacing alloy Stellite 1 on AISI-SAE 4340 steel was prevented through locally preheating the substrate prior to the deposition process. The result showed that the preheated sample revealed more uniform structure and was less prone to cracking. Fallahi et al. [21] introduced an efficient approach to reduce the residual stress in the welding process by a preheating treatment. The residual stress was decreased when the preheating temperature was increased from 25 °C to 400 °C. When the coating is prepared at a high circumstance temperature, the residual stress in the coating is released and the phase transformation possibly happens. These changes will result in the elimination in strain-hardening effect and the evolution in coating microstructure. As a result, wear resistance of the coating will be inevitably influenced at the same time. Therefore, it is necessary to study the influence of preheating and postheat treatments on wear resistance of the coating with a low residual stress. The investigation will contribute to confirming the appropriate treatment temperature, in which not only the cracking susceptibility is reduced, but also excellent wear resistance is maintained. However, there are hardly investigations about the issue. In this study, the Ti2Ni/TiNi based coatings were prepared on Ti6Al4V at different circumstance temperatures (25 °C, 400 °C, 600 °C and 800 °C) by laser cladding, then were preserved for 3 h. The relationship between cracking susceptibility and wear resistance of the coatings was investigated in detail. Finally, an appropriate temperature was determined to obtain the best comprehensive mechanical properties of the coatings.
2. Experimental procedures Ti6Al4V [composition (in wt.%): 6.5 Al, 4.26 V, 0.1 C, balance Ti] which was cut into the samples with a dimension of 30 mm ⁄ 20
mm ⁄ 10 mm was used as the substrate. F102 Ni-based alloy powder [composition (in wt.%): 75 Ni, 1 C, 16 Cr, 3.5 B, 4.5 Si] was selected as the cladding material. The substrate surface was brushed with a binder (4% polyvinyl alcohol) and then placed in a cubic model with a dimension of 30.2 mm ⁄ 20.2 mm ⁄ 10.8 mm. A space with a precise height of 0.8 mm was formed above the substrate. The powder was filled into this the space and compacted to form a preplaced layer at 30 MPa for 3 mins using a tablet machine. This method not only ensures the uniformity in thickness of the preplaced layer, but also reduces its porosity. Before laser cladding, the samples were preheated to 400 °C, 600 °C and 800 °C, and were maintained for half an hour in a resistance furnace. After that, the furnace lid was opened and laser cladding was directly carried out on the samples placed in the furnace. The circumstance temperature during laser cladding was nearly maintained at the initially set temperature (400 °C, 600 °C, and 800 °C), because the laser cladding time was very short (about 15 s) and the furnace was still in work during laser cladding. Laser cladding was performed using an YLS-5000 fiber laser processing system with an applied power of 3 KW, a spot diameter of 6 mm, and a scanning speed of 10 mm/s. The parameter was optimized by ensuring a good interface fusion and a smooth coating surface. The sample without preheating (25 °C) as the contrast was prepared at the same processing parameter, and the obtained coating was named as coating I. The coatings prepared at 400 °C, 600 °C and 800 °C were named as coatings II, III and IV. After laser cladding, the samples were preserved at 400 °C, 600 °C and 800 °C for 3 hours, respectively. Phase constituents of the coatings were analyzed through a PANalytical X’ Pert Pro X-ray diffractometer (XRD). Macromorphologies of the coatings were observed with a VHX-600 K optical microscope (OM). Microstructural characterization was conducted on the cross-sections of coatings using an S-3400 scanning electron microscope (SEM) coupled with a GENESIS EDAX energy dispersive spectrometer (EDS). Microhardness of the cross-sections of coatings was measured by a HXD-1000TMSC/LCD microhardness tester with a load of 200 gf applied for 15 s. Nanoindentation tests were performed using a TriboIndenterÒ (Hysitron Corporation, USA) with a diamond Berkovich indenter. The indenter was also used as an AFM tip to imaging the indented surface after the tests. A peak load of 10 mN was selected in the tests, and the loading and unloading times were set at 5 seconds each and the maximum load was also maintained for 5 seconds. The first nanoindentation was obtained at the zone of 0.1 mm distance from the coating surface. The next indentations were prepared at different zones with the same distance interval (0.4 mm). Three zones in every coating were selected for the tests. Room-temperature dry friction tests were performed to evaluate wear resistance of the coatings using a CFT-I ultra-functional wear testing machine in the ball-on-disc reciprocating motion mode. The sliding time was 180 min, the applied load was 30 N, and the reciprocating sliding speed was 5 m/min. YG6 balls with a diameter of 6 mm were used as the counterparts. Wear surface morphologies of the coatings were observed with SEM. Wear volumes of the coatings were measured using a wear profiler.
3. Results and discussion 3.1. Macromorphologies Fig. 1 shows the macromorphologies of the coatings prepared at different circumstance temperatures. The profiles of the coatings are very similar, presenting the convex shape. No obvious cracks and unmelted particles are observed. A good metallurgical bonding is formed between coatings and substrate. The width (6.21 mm at
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Fig. 1. Morphologies of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
25 °C, 7.00 mm at 400 °C, 7.08 mm at 600 °C and 7.20 mm at 800 °C) and thickness (1.98 mm at 25 °C, 2.10 mm at 400 °C, 2.44 mm at 600 °C and 2.84 mm at 800 °C) of the four coatings present an increase tendency with the increase in circumstance temperature. Dilution rate g can be calculated as follows [22]:
g¼
S1 100% S1 þ S2
ð1Þ
where S1 is the molten area of the substrate, and S2 is the area of the coating above the surface of the substrate. The dilution rate is determined to be 64.40% (25 °C), 70.21% (400 °C), 72.53% (600 °C) and 74.29% (800 °C). The dilution rate of the coatings is increased with the increase in circumstance temperature due to more extra energy provided in the cladding system. 3.2. Microstructural characterization Fig. 2 illustrates the XRD spectra of the coatings prepared at different circumstance temperatures. The spectrum of the coating prepared at 25 °C is very simple. Main diffraction peaks are in accordance with the Joint Committee on Powder Diffraction Standards (JCPDS) cards (No. 03-065-5537 for TiNi, No. 00-018-0898 for Ti2Ni, No. 03-065-8698 for TiB2, and No. 03-065-8804 for TiC). The three strongest lines of every phase are in good agreement with those in the JCPDS cards. Therefore, the coating is composed of TiNi, Ti2Ni, TiB2, and TiC. When the preheating temperature is increased to 400 °C, two extra diffraction peaks at 2h = 39.1° and 2h = 39.6° are observed and the intensity of a diffraction peak at 2h = 42.7° is significantly increased. It can be confirmed that a new phase of Cr2Ti was in situ synthesized in the coating prepared at 400 °C. The XRD spectrum of the coating prepared at 600 °C is highly similar to that of the coating prepared at 400 °C, indicating that their phase constituents are same. As the preheating temperature was further increased to 800 °C, a significant change is visible. Three new diffraction peaks appear at 2h = 47.8°, 2h = 71.0° and 2h = 76.7°. In terms of the JCPDS cards, the other new phase of Ti3Al was formed in the coating prepared
at 800 °C. Other than that, the intensities of two diffraction peaks (2h = 41.5° and 2h = 42.7°) also present the increasing tendency when compared with those in Fig. 2b and c. Considering that the contents of Ni and C are same in the four coatings, the increase in intensity of the two diffraction peaks should be attributed to the increase in volume fraction of Ti2Ni and Cr2Ti. As mentioned above, the dilution rate of the coatings is increased with the increase in circumstance temperature (64.40% at 25 °C, 70.21% at 400 °C, 72.53% at 600 °C, and 74.29% at 800 °C). The increase in dilution rate means that more Ti enters into the molten pool during laser cladding. As a result, some compounds rich in Ti (Cr2Ti and Ti3Al) are formed successively with the increasing in circumstance temperature, and the volume fractions of three compounds rich in Ti (Ti2Ni, Cr2Ti and Ti3Al) also show the increasing tendency. Fig. 3 shows the Backscattered Electron (BSE) images of representative microstructures of the four coatings. A large number of acicula-shaped structures are distributed uniformly in the matrix of coating I. As indicated in the XRD spectra, the reinforcements in the four coatings mainly include TiB2 and TiC. When the circumstance temperature is increased to 400 °C and 600 °C, the microstructure of coatings II, III is similar to that in coating I. However, the density of acicula-shaped structures exhibits the downward trend. With the increase in circumstance temperature (800 °C), the density of the reinforcements is further decreased. Moreover, the morphology of the reinforcements is transformed from coarse acicula shape into fine feather shape. It is well known that the growth of the reinforcements strongly depends on the atomic content in the molten pool. The dilution rate of coating III is highest among the four coatings, which indicates that the concentration of B and C is also lowest in the molten pool. During solidification of the molten pool, the growth of the reinforcements in coating III will be restricted greatly when compared with the other three coatings. As a result, the reinforcements in the coating prepared at 800 °C are finer. The microstructure in a high-magnification image further reveals that coating I is composed of six phases (shown in Fig. 4), corresponding to cupped light gray phase (phase 1), protuberant gray white phase (phase 2), grey strip-shaped/spherical particles (phase 3), block acicula-shaped particles (phase 4), rodlike parti-
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Fig. 2. X-ray diffraction spectra of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
Fig. 3. Microstructures of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
cles (phase 5), and equiaxed particles (phase 6). Phases 1 and 2 with a high volume fraction can be regarded as the matrix, phases 3, 4, 5 and 6 should be the in situ synthesized reinforcements. The rodlike (phase 5) and equiaxed particles (phase 6) as the primary phase exist independently in the matrix. However, the comparatively fine grey strip-shaped/spherical particles (phase 3) are uniformly distributed in the block acicula-shaped particles (phase 4), which can be identified as a eutectic structure. The compositions of those phases with different morphologies were analyzed by EDS, and the results are shown in Table 1. Phases 1 and 2 are rich in titanium and nickel, accompanied with a small amount of Cr, Al and Si, which should be two kinds of Ti-Ni compounds. Ti
and Ni atoms are easily replaced by (Cr and Al) and Si atoms, respectively due to their similar atomic radii and electronegativities. The atomic percentage of Ti, Cr and Al in phase 1 is 48.00 at. %, and that of Ni and Si is 49.70 at.%. Thus the cupped light gray phase (phase 1) can be identified as the TiNi solid solution. Similarly, the protuberant gray white phase (phase 2) can be regarded as Ti2Ni since the atomic ratio of (Ti + Al + Cr) and (Ni + Si) is about 1.54:1. The grey strip-shaped/spherical particles (phase 3), rodlike particles (phase 5) and equiaxed particles (phase 6) with very similar chemical compositions are rich in titanium and carbon. The block acicula-shaped particles (phase 4) mainly consists of titanium (15.37 at.%) and boron (84.08 at.%). Combined with the
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Fig. 4. Typical microstructure of coating I.
Table 1 EDS analyses (in at.%). Coatings
Phase
Ti
C
B
Ni
Cr
Al
Si
Phase information
I
1 2 3 4 5 6
32.49 39.06 48.34 15.37 40.68 51.65
1.24 1.54 51.66 0.55 59.32 48.35
– – – 84.08 – –
48.12 26.24 – – – –
4.24 13.66 – – – –
11.27 0.71 – – – –
1.58 8.46 – – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
II
1 2 3 4 5 6
41.01 34.74 61.72 13.56 41.52 50.63
2.16 2.18 33.54 – 58.48 49.37
– – – 86.44 – –
42.84 27.91 2.36 – – –
4.80 17.49 0.46 – – –
6.58 5.04 0.52 – – –
1.26 8.85 0.84 – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
III
1 2 3 4 5 6
40.46 36.23 60.30 14.05 45.92 48.69
1.13 2.14 39.70 – 54.08 51.31
– – – 85.95 – –
45.29 25.50 – – – –
4.10 16.04 – – – –
6.85 5.45 – – – –
1.00 10.47 – – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
IV
1 2 3 4 5 6
51.07 41.93 51.58 22.41 44.35 47.33
2.81 3.18 48.42 – 55.65 52.67
– – – 77.59 – –
26.86 26.11 – – – –
4.80 10.23 – – – –
8.73 9.70 – – – –
4.29 5.86 – – – –
TiNi Ti2Ni TiC TiB2 TiC TiC
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XRD analyses, phases 3, 5, 6 can be identified as TiC and phase 4 should be TiB2. The compositions of six phases in the other coatings are very similar to those in coating I (shown in Table 1). The reinforcements in the composite coatings are TiC and TiB2. The microstructural evolution of the coating can be interpreted through a TiB2-TiC phase diagram [23]. The eutectic structure is formed when the mole fraction of TiC is 0.2175. The mole fraction of TiC in the molten pool can be calculated as about 0.3432 in this study, which indicates that the content of TiC is located in the hypereutectic area. Therefore, TiC as the primary phase is first precipitated from the liquid and then grows into the rodlike or equiaxed particles. TiC has a B1 crystal structure and the geometry of TiC crystal cell presents the octahedral structure. The TiC crystal is inclined to grow along the h100i directions preferentially due to its face centered cubic structure [24]. When the degree of supercooling in the molten pool is small, the octahedral TiC crystal uniformly branches in the six h100i directions with the similarly same growth speed, resulting in the formation of equiaxed TiC particles. However, the crystal in a specific direction may grow faster than the other directions during subsequent solidification because of the differences in temperature and concentration distributions in different directions [25]. As a result, the rodlike TiC particles are formed. With the decrease in temperature in the molten pool, a eutectic structure composed of TiB2 and TiC is formed by heterogeneous nucleation and grows into the block acicula-shaped structure [26]. Afterwards, Ti2Ni/TiNi are formed.
3.3. Nanoindentation tests Fig. 5 shows the load-depth curves of the coatings prepared at different circumstance temperatures. According some data obtained from Fig. 5, some mechanical properties (such as the residual stress) of the coatings can be extracted by the OliverPharr method. The key parameter in the calculation is the contact area, which is directly influenced by the ratio of the final indentation depth (hf) and the maximum indentation depth (hmax). The ratio reflects the degree of pile-up deformation in the coating. A serious pile-up deformation will generate around the indentation edge when hf/hmax is close to 1. As a result, the area deduced from the load-depth curves is far less than the true area by as much as 60% [27]. This leads to a serious overestimation of mechanical properties of the coating, even a significant error (exceeding fifty percent) may be produced. When hf/hmax < 0.7, very little pile-up occurs and a comparatively accurate result can be obtained. In this study, the values of hf/hmax at different zones of the coatings are listed in Table 2. It is clear that all values are lower than 0.7. Fig. 6 shows the two-dimensional and three-dimensional morphologies of a typical nanoindentation of coating III. No obvious pile-up deformation is observed in the zone around the nanoindentation. The Suresh model is commonly used to distinguish and calculate the residual stress of the material, which is mainly based on the difference in maximum depth between a stress-free material
Fig. 5. Load-depth curves of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
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Table 3 Residual stress at different zones in the coatings (GPa).
Coatings
0.1 (mm)
0.5 (mm)
0.9 (mm)
Reference
Coatings
0.1 (mm)
0.5 (mm)
0.9 (mm)
I II III IV
0.61 0.58 0.59 0.56
0.50 0.62 0.60 0.58
0.57 0.57 0.59 0.58
Coating IV
I II III IV
2.51 1.72 1.03 0.00
2.56 1.94 1.05 0.00
3.27 0.58 0.52 0.00
Coating III
I II III IV
1.78 0.81 0.00 1.24
1.60 1.10 0.00 1.31
1.96 0.04 0.00 0.49
Coating II
I II III IV
1.13 0.00 0.92 2.34
0.32 0.00 1.39 3.05
2.00 0.00 0.04 0.53
Coating I
I II III IV
0.00 1.41 2.51 4.26
0.00 0.28 1.74 3.48
0.00 2.44 2.30 1.67
and a material with residual stress. Therefore, the selection of a stress-free material is a very key factor for the calculation of the residual stress. Considering that the microstructure of the laser cladding coating is very complex, it is difficult to prepare a stress-free coating with similar microstructure by the other methods. A feasible method is to select a material or a specific zone in a material with a comparatively low stress as the stress-free reference among all materials, which had been used widely in previous studies. Zhu et al. [27] studied the residual stress in the plasmasprayed FeCrBSi coating by the nanoindentation method. As the residual stress in the interface between FeCrBSi coating and Ni/Al transition layer was close to zero, the interface was regarded as the stress-free reference. Zhang et al. [28] investigated the residual stress in 304 stainless steel by the nanoindentation method. The annealed stainless steel was selected as the stress-free sample to measure the residual stress in the unannealed sample. The residual stress was about 381 MPa, which was very close to the result tested by the XRD method (350 ± 23 MPa). Wang et al. [29] measured the residual stress in the plasma sprayed Fe-based coating using the nanoindentation method. A coating with the 0.5 mm thick surface removed by the wire-electrode cutting method was selected as the stress-free coating since the residual stress was released fully. According to the fundamental theory of materials science, the residual stress in the coating can be reduced by reducing the temperature difference between environment and cladding. Therefore, the coating prepared at 800 °C was selected as the stress-free coating in our study, In order to verify the rationality, the residual stress in the four coatings was calculated when coatings I, II, III and IV were selected as the stress-free coating, respectively. As shown in Table 3, the residual compressive stress exists in some coatings when coatings I, II and III are selected as the stress-free coating. The residual tensile stress exists in the three coatings when coating IV is regarded as the stress-free coating. Some studies had confirmed that the residual tensile stress was produced easily in the laser cladding coating due to the rapid
heating and cooling process during laser cladding [29,30]. Therefore, it is appropriate to select coating IV as the stress-free coating. The residual stresses in the other coatings were calculated by the following equations [31]: For the tensile residual stress:
r¼H 1
2
h0 h
! ð2Þ
2
For the compressive residual stress:
r¼
2
H h0 1 sin a h2
! ð3Þ
where h and h0 are the penetration depth in the coatings with and without the residual stress, respectively, H is the hardness of the coatings, and a is related to the indentation angle of the indenter (a = 24.7 for a Berkovich indenter). Fig. 7 indicates the load-depth curves obtained at different zones of the four coatings. The maximum depth in coating I always reaches the maximum value. At the zones with 0.1 mm and 0.5 mm distance from the surfaces of the coatings, the maximum depth presents the decreasing trend with the increase in circumstance temperature (shown in Fig. 7a and b). At the zones with
Fig. 6. Two-dimensional and three-dimensional morphologies of a typical nanoindentation of coating III.
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Fig. 7. Load-depth curves for the coatings: (a) 0.1 mm, (b) 0.5 mm and (c) 0.9 mm.
about 0.9 mm distance from the surfaces of coatings I, II and III (shown in Fig. 7c), the maximum depth is approximately same. When the same load is applied to the coatings, the different indentation depth is produced, which should be attributed to the existence of the residual stress in the coatings. The tensile stress tends to facilitate the indentation process. However, the compressive stress will produce the contrary effect in the indentation formation. Therefore, it is concluded that the tensile stress exists in coatings I, II, and III when coating IV is regarded as the stressfree reference. The residual tensile stress in the coatings was calculated by the Suresh model (shown in Table 3). The residual tensile stress at the zones with different distance from the surface of coating I always reaches the maximum value (2.51 GPa in 0.1 mm, 2.56 GPa in 0.5 mm and 3.27 GPa in 0.9 mm). At the zones with 0.1 mm distance from the surfaces of the coatings, the residual tensile stress is reduced about 25.1% (1.72 GPa) and 40.3% (1.03 GPa) with increasing the circumstance temperature to 400 °C and 600 °C. At the zones with 0.5 mm distance from the surfaces of the coatings, the residual tensile stress in the coatings is reduced about 24.4% (1.94 GPa for 400 °C) and 45.6% (1.05 GPa for 600 °C). However, the residual tensile stress in coatings II and III approaches zero at the zones with 0.9 mm distance from the surfaces of the coatings. It can be concluded that the residual stress in the coatings prepared at high circumstance temperatures could be reduced obviously in whole. Fig. 8 shows the microhardness distribution throughout the coatings. The microhardness shows a downward trend along the
depth direction of the cross-sections of the coatings. Considering the change in microhardness, three regions can be clearly identified, corresponding to the coating, the transition zone and the substrate. The average microhardness values of coating I at the three zones are 1062.2 HV0.2, 624.2 HV0.2, and 350.8 HV0.2. The average microhardness value of the coating is about 3 times that of the sub-
Fig. 8. Microhardness distribution across the cross sections of the coatings.
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strate due to the uniform distribution of a large number of highhardness TiC (30 GPa) [32] and TiB2 (34 GPa) [33] particles within TiNi and Ti2Ni intermetallic compounds. The existence of the transition zone contributes to forming a good metallurgical bonding between coating and substrate. Moreover, the transition zone plays a buffer effect between hard coating and soft substrate. The coating and the substrate have different elasticity modulus and thermal expansion coefficient, which leads to different deformation when the coating is subject to the external force or temperature variation in service. As a result, the huge stress concentration will generate in the junction interface between coating and substrate. The crack will initiate and propagate when the stress exceeds the bonding strength. Therefore, the transition zone can effectively reduce the cracking sensibility of the coating. The average microhardness values of coatings II, III and IV are 1031.2 HV0.2, 1028.2 HV0.2, and 954.7 HV0.2, respectively. It is clear that the average microhardness values of the coatings are reduced as increasing the circumstance temperature. The decline in microhardness and residual stress of the coatings has the close relationship with the changes in microstructure (especially atomic migration). The high residual stress usually exists in the coating fabricated by laser cladding due to its rapid heating and cooling characteristics. The strainhardening effect resulting from the residual stress will further improve the microhardness of the coating. The number of defects in the coating is reduced due to the rapid atomic migration when the coating is prepared and maintained at high circumstance temperature. Accompanied with this, the residual stress is released and the microhardness is decreased accordingly. On the other hand, the reduction in reinforcement density can also cause the decrease in microhardness to a certain extent. 3.4. Wear resistance Wear profiles of the coatings are shown in Fig. 9. The wear volume is 0.1912 mm3 for coating I. The wear volumes are increased by nearly 47.9% (0.2828 mm3 for coating II), 95.2% (0.3732 mm3 for coating III) and 220% (0.6073 mm3 for coating IV) with the increase in circumstance temperature. Fig. 10 shows the surface morphologies of coatings I, II, III and IV after the wear tests. Many spalling holes are observed on the worn surface of coating I (shown in Fig. 10a). The other zones except the spalling holes are comparatively smooth. Coatings II, III and IV present the similar worn surface morphology (shown in Fig. 10b–d). However, the number of the spalling hole on the worn surfaces of the coatings is decreased with increasing the circumstance tem-
Fig. 9. Wear profiles of the coatings.
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perature. The distribution and morphological characteristic of the hole resulting from the spalling of wear debris can be used to reveal the formation mechanism of the debris. The spalling hole generally generates in the center or edge of the reinforcements and hardly generates in the matrix. During sliding friction, the difference in elasticity modulus between reinforcement and matrix leads to their different deformation when a load is applied to the contact surface. The stress concentration will generate in the interface. The crack may initiate and propagate along the interface, finally into the matrix when the stress exceeds the interface bonding strength. Moreover, the crack may generate within the reinforcement and extend toward the matrix due to its low toughness. The bottom of the spalling hole is actually the fracture surface, which records the details of crack propagation. The riverlike pattern can be clearly observed at the bottom of the spalling hole, which presents the typical brittle-fracture characteristic. The images with the high magnification clearly indicate that the morphology of the fracture surface of coating I well accords with that of the typical intergranular fracture (shown in Fig. 11a). Many grains can be clearly observed, and the whole surface is comparatively smooth due to the insignificant change in contrast. A small number of white tearing edges can be only visible in the grain boundaries. It can be concluded that the plastic deformation does not almost occur during crack propagation. That is to say, the crack propagates with less resistance. The large residual tensile stress in coating I will further promote the crack initiation and propagation. The impurity elements (such as O, N) have no enough time to escape from the molten pool during laser cladding due to its rapid heating and cooling characteristic. They may aggregate in the grain boundaries of TiNi and Ti2Ni, which reduce the bonding strength of the boundaries. The crack tends to initiate and propagate along the grain boundaries during wear, resulting in the intergranular fracture. Therefore, wear of coating I is mainly dominated by brittledebonding. As shown in Fig. 11b, the surface of the spalling hole in coating III is rougher when compared with that of coating I, which indicates that the comparatively large plastic deformation occurs during crack propagation. The plastic deformation will consume the extra energy and hinder the crack propagation. On the other hand, the decrease in residual tensile stress will further restrain the crack propagation. Therefore, the brittle-debonding is relieved with increasing the circumstance temperature. Moreover, some slight furrows and scratches are clearly observed on the worn surface of coating III (as shown in Fig. 11b). It can be concluded that wear mechanism of the coatings is transformed from the single brittle-debonding into the combination of the micro-cutting and the brittle-debonding when the processing temperature is changed from room temperature to high temperature. The phenomenon is closely related to the changes in hardness and residual stress of the coatings. As analyzed above, the debris is mainly formed by two mechanisms corresponding to micro-cutting and brittle-debonding. The resistance to micro-cutting mainly depends on hardness of the coating, the resistance to brittle-debonding is closely related to plasticity and residual stress of the coating. With the change in preheating temperature, the resistance to micro-cutting and brittle-debonding of the coating will present the regular change. Among the four coatings, the hardness of coating I is comparatively high due to the high density of the reinforcements and the strong strain-hardening effect, indicating that the coating has the powerful resistance to micro-cutting. The coating with high hardness usually possesses the comparatively low plasticity. Moreover, a very high residual tensile stress exists in coating I without preheating. The mutual effect of the two factors will significantly accelerate the crack initiation and propagation in coating I when it is in service. That is to say, coating I exhibits very high cracking susceptibility and very less resistance to brittle-debonding. When the
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Fig. 10. Morphologies of the wear surfaces of the coatings: (a) 25 °C, (b) 400 °C, (c) 600 °C and (d) 800 °C.
Fig. 11. Typical morphologies of the wear surface of the coatings: (a) 25 °C and (b) 600 °C.
preheating temperature is increased to 400 °C and 600 °C, the hardness of the coating is reduced from HV 10620.2 to about 1030 HV0.2 due to the reduction in density of reinforcements and the weakening in strain-hardening. It indicates that the two coatings present the weaker resistance to micro-cutting than coating I. However, their resistance to brittle-debonding is improved to a certain extent along with the increase in plasticity and the decrease in residual stress. The combination of the two effects (micro-cutting and brittle-debonding) results in about 0.3280 mm3 loss of coatings II and III, which is higher than that of coating I. Similarly, when the preheating temperature is increased to 800 °C, the resistance to micro-cutting of coating IV is further reduced, accompanied with the increase in resistance to brittle-debonding.
The wear volume of coating IV is further enhanced to 0.6073 mm3. According the change in wear volume with the increase in preheating temperature, it can be concluded that the microcutting plays the important role in debris formation. The debris is immediately produced with the movement of the counterpart during the micro-cutting process, while the brittle-debonding process involves crack initiation, propagation and debonding, which needs more time. Thus, the micro-cutting is easier to produce debris, when compared with the brittle-debonding. Based the above analyses, it can be concluded that laser cladding at the high temperature is useful to reduce cracking susceptibility in the coatings. However, wear resistance of the coatings will be decreased to a certain extent. When the preheating temperature
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exceeds 600 °C, the wear volume of the coating (800 °C) is increased about 220% when compared with that without preheating. Wear resistance of the coating is weakened significantly, which greatly reduces its service life. When the preheating temperature is less than 600 °C, the wear volume of the coating (400 °C) is only increased about 47.9% when compared with that without preheating. The coating still presents the excellent wear resistance. However, it also has a high cracking susceptibility since a comparatively high residual tensile stress exists. Comparatively speaking, the wear volume of the coating prepared at 600 °C is slightly lower than that of the coating prepared at 400 °C. However, the residual stress is further reduced about 40%. Therefore, the suitable preheating temperature should be 600 °C, at which the comparatively high wear resistance is maintained on the premise that the residual stress is relieved. 4. Conclusions (1) TiNi/Ti2Ni-based composite reinforced by TiC and TiB2 were prepared on the Ti6Al4V substrate at different circumstance temperatures (25 °C, 400 °C, 600 °C and 800 °C) by laser cladding, then were preserved for 3 h. (2) The tensile stress existed in the coatings when the coating prepared at 800 °C was regarded as the stress-free reference. The average residual stress in the surface of coating prepared at 25 °C reached the largest value of about 2.79 GPa and presented the decreasing tendency with increasing the circumstance temperature (1.03 GPa at 400 °C, 0.52 GPa at 600 °C and 0 GPa at 800 °C). (3) The wear volumes of the coatings prepared at 25 °C, 400 °C, 600 °C and 800 °C were increased with the increase in circumstance temperature (0.1912 mm3 at 25 °C, 0.2828 mm3 at 400 °C, 0.3732 mm3 at 600 °C and 0.6073 mm3 at 800 °C). The brittle-debonding was predominant in the debris formation of the coating prepared at 25 °C. Wear mechanism was transformed into the combination of the microcutting and brittle-debonding with the increase in processing temperature. The micro-cutting caused a more serious wear loss of the coating, when compared with the brittledebonding. (4) The processing temperature should be 600 °C, at which not only the high wear resistance was maintained, but also the cracking susceptibility was reduced.
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