稀有金属(英文版) 2020,39(12),1402-1412

Interface characteristic and mechanical performance of TiAl/Ti₂AlNb diffusion bonding joint with pure Ti interlayer

Lei Zhu Bin Tang Ming-Xuan Ding Yan Liu Xiao-Fei Chen Shao-Peng Yan Jin-Shan Li

State Key Laboratory of Solidification Processing,Northwestern Polytechnical University

Shaanxi Key Laboratory of High-Performance Precision Forming Technology and Equipment,Northwestern Polytechnical University

作者简介：*Bin Tang,e-mail:toby@nwpu.edu.cn;

收稿日期：25 March 2020

基金：financially supported by the National Natural Science Foundation of China (No.51771150);the National Key Research and Development Program of China (No. 2016YFB0701303);the Aeronautical Science Foundation of China (No.201936053001);the Research Fund of the State Key Laboratory of Solidification (NWPU),China (No.2019-TS-07);

Interface characteristic and mechanical performance of TiAl/Ti₂AlNb diffusion bonding joint with pure Ti interlayer

Lei Zhu Bin Tang Ming-Xuan Ding Yan Liu Xiao-Fei Chen Shao-Peng Yan Jin-Shan Li

State Key Laboratory of Solidification Processing,Northwestern Polytechnical University

Shaanxi Key Laboratory of High-Performance Precision Forming Technology and Equipment,Northwestern Polytechnical University

Abstract：

Solid-state diffusion bonding(DB) of TiAl alloy and Ti₂ AlNb alloy was carried out using pure Ti as an interlayer at 1000℃ under 20 MPa for 60-120 min.The effects of bonding times on the interfacial microstructure and mechanical performance of the TiAl/Ti/Ti₂ AlNb bonded joints at room temperature(RT) were investigated detailly.The results demonstrated that the diffusion layers(DLs) mainly consisted of four characteristic layers,(Ⅰ) single coarse α₂ phase adjacent TiAl alloy,(Ⅱ) single refined α₂ phase at the bonding interface,(Ⅲ) equiaxed/acicular α₂ phase embedded in β phase adjacent Ti₂ AtNb alloy and(IV) both equiaxed α₂ phase and acicular O phase embedded in β phase adj acent Ti₂ AlNb alloy,respectively.The thickness of the four layers increased with the increasing of the bonding time.The growth of DLs is controlled by diffusion and the reaction rate constant k for region Ⅰ,Ⅱ,Ⅲ and Ⅳ are 1.22×10^-6,1.27×10^-6,2.6×10^-7 and 7.7×10^-7 m·s^-1/2,respectively.Meanwhile,the interface α₂ grain grows up without texture.The maximum tensile strength of 281 MPa was maintained at1000℃ for 90 min under the pressure of 20 MPa.Consequently,the phase transformation and dynamic recrystallization behavior of the DLs were discussed.

Keyword：

Sputtering; Bonding interface; Microstructure; Growth kinetics; Mechanical performance;

Received： 25 March 2020

1 Introduction

As one of pivotal high-temperature structure engineering material,TiAl-based alloy has been used in aerospace and auto industry for several years due to its attractive"strong and light"properties,such as excellent high-temperature strength,low density,good creep and outstanding oxidation resistance [ 1, 2, 3, 4, 5, 6, 7] .Nevertheless,it is regretful that the engineering application of TiAl-based alloy is also limited due to the intrinsic brittleness and poor hot workability [ 8, 9] .Therefore,it is necessary to using welding technology to expand the practical engineering application of TiAl-based alloy.In recent years,some melt weld methods have been applied to join the TiAl-based alloy [ 10, 11, 12, 13] .Compared with the melt weld methods,vacuum solid-state diffusion bonding (DB) has been thought to be an appropriate way to fabricate TiAl-based alloy [ 14, 15, 16, 17] ,which can a^void the micro-cracks and reduce the thermal stress at the bonding interface.Whereas,the brittle intermetallic compounds (IMCs),AlNb₂ phase or Al(Nb,Ti)₂ structure phase [ 18, 19] ,are always formed at the interface of the joint yet when using the DB [ 20, 21, 22] .

On the other hand,there are still many difficulties for direct diffusion bonding dissimilar intermetallic compounds on account of the different thermal expansion coefficients and chemical compositions.Accordingly,it is difficult to bond the dissimilar intermetallic compounds without defects and internal stresses.Therefore,to inhibit the formation of harmful IMCs and reduce the residual stress at the bonding interface,the addition of a proper interlayer was conducted to obtain a resultant joint [ 23, 24, 25, 26] .In the present work,the pure Ti interlayer was selected to join the TiAl alloy and Ti₂AlNb alloy using vacuum solidstate diffusion bonding.TiAl alloy was suitable for fabricating the blades because of its excellent creep resistance at high-temperature (700-850℃).Although the working temperature (650-750℃) of Ti₂AlNb alloy is lower than that of TiAl alloy,Ti₂AlNb alloy possesses relatively higher ambient ductility and crack propagation resistance than TiAl alloy.It is meaningful to realize weight reduction by bonding the two intermetallic compounds.In addition,the pure Ti was chosen as an interlayer in this work was mainly because of the following reasons:(1) Ti is the primary component of two base materials,avoiding the IMCs formed at the bonding interface,which has outstanding suitability with two as-received alloys.(2) the contact between TiAl alloy and Ti₂AlNb alloy is compact on account of the addition of pure Ti,which promotes the elimination of micro-voids.(3) activity of Ti atom is high during the vacuum solid-state diffusion bonding,meanwhile,there is a significant number of grain boundaries within the interlayer formed by direct current (DC) magnetron sputtering,leading to a sufficient diffusion across the bonding interface.

The objective of this present work is to investigate the influences of bonding time on the interfacial microstructures and the mechanical properties of the TiAl/Ti/Ti₂AlNb bonding joints.Besides,the feasibility of diffusion bonding of TiAl alloy and Ti₂AlNb alloy using pure Ti as an interlayer,which is essential for enlarging their engineering applications,was also evaluated.The relationship between the microstructural evolution and the mechanical performance of the bonding joints at room temperature (RT) was investigated.Moreover,the dynamic recrystallization and phase transformation behavior were discussed.Finally,diffusion bonding mechanism was analyzed based on the microstructural observation.

2 Experimental

The ingot of TiAl alloy was manufactured by vacuum arc remelting (VAR) three times following with hot isostatic pressing (HIP) at 1280℃for 4 h under the pressure of140 MPa.After HIP,the ingot was forged at 1170℃to obtain a forging pancake with a size ofΦ420 mm×100mm.The initial microstructure with equiaxed structure is displayed in Fig.1a,including theβ/B2 phase (bright),α₂phase (gray) andγphase (dark),as given in Fig.1c.The actual chemical composition is shown in Table 1.

The Ti2AlNb alloy with a nominal composition of Ti-22Al-25Nb (at%) used in this work was fabricated by triple VAR,and then,the ingot was free forged by a multi-pass forging process.The size of the forging bar wasΦ220 mm×500 mm finally.The forging microstructure of the Ti2AlNb alloy is shown in Fig.1b.The microstructure consists of gray rim O phase around the darkest ovalα2 phase and the gray acicular O phase enmeshed in the brightestβ/B2 phase.Figure 1d also demonstrates X-ray diffraction (XRD) pattern of Ti2AlNb,including O phase,α2 phase andβ/B2 phase.The detailed chemical composition is also shown in Table 1.

The forging TiAl alloy and Ti2AlNb alloy were machined by electric-discharge machining into cuboids with the dimensions of 50 mm×30 mm×20 mm.The faying surface was grounded subtly followed by mechanical polished using 1.5 mesh abrasive paste,and the asreceived specimens were ultrasonic cleaned in alcohol for～5 min to clean the faying surface.Finally,the samples were vacuum encapsulated for diffusion bonding.The interlayer of pure Ti was deposited by DC magnetron sputtering.The total thickness of the interlayer was about10μm,which was composed of two 5μm layers sputtering on the faying surface of two substrates.

Diffusion bonding was taken place in a vacuum diffusion bonding machine (ZC-ZK YL40) at 1000℃and20 MPa for different bonding time of 60-120 min.The selection of temperature and time were based on the previous research results [ 18] .The heating rate was set as10℃·min^-1 from room temperature to target temperature and furnace cooling (FC) to the room temperature during the diffusion bonding process,as shown in Fig.2.The pressure of 20 MPa was implemented when the bonding temperature reached until furnace cooling.After the bonding process,the samples were fabricated by the conventional metallographic procedure across the bonding interface.

The microstructure of the interface was characterized by scanning electron microscope (SEM,TESCAN VEGA3)with electron back-scattered diffraction (EBSD),and the dispersion of elements in the diffusion layers (DLs) and metallic compounds was measured by electron probe micro-analyzer (EPMA,JEOL JXA-8230).Notably,the schematic diagram of the coordinate system in EBSDanalysis is shown in Fig.3a.The specimens for EBSDwere produced by the once mechanical grinding and electropolishing on a Struers Lectropol-5 (Denmark,Struers Corporation) in a solution of 15 vol%n-butyl alcohol,5vol%perchloric acid and 80 vol%methanol with a voltage of 35 V holding for 30 s.EBSD analytical test was executed with a scan step size of 0.2μm to obtain the detailed orientation information of the interface,the geometrically necessary dislocations (GND) maps from EBSD data were calculated by using the software named HKL Channel 5.Transmission electron microscopy (TEM) measurements were conducted at Themis Z Double Cs Corrector microscope operated at 300 kV.TEM specimens were preparedby focused ion/electron double beam electron microscopy(FIB).Moreover,to assess the joints quality,a nonstandard tensile test was designed to appraise the tensile properties of the joints in an electronic universal testing machine controlled by microcomputer (SUNS) with an initial strain rate 1×10^-4 s^-1 at RT.Figure 3b shows the schematic diagram of nonstandard tensile specimens whose bonding interface is situated in the middle of the tensile samples.

Fig.1 SEM images of based materials in BSE mode:a TiAl alloy and b Ti2AlNb alloy;XRD patterns of c TiAl alloy and d Ti2AlNb alloy

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Table 1 Chemical compositions of TiAl and Ti2AlNb alloys (at%)

Fig.2 Curve of diffusion bonding process

3 Results and discussion

diffusion behavior of bonding interface

3.1 Microstructure characteristic and elemental

Figure 4a,b shows the typical micros true ture of TiAl/Ti2AlNb bonding interface without pure Ti interlayer andwith pure Ti interlayer at 1000°C-20 MPa-90 min,respectively.It is self-evident that there exists a discon?tinuous equiaxed white grain produced at the bonded interface without an interlayer,which is similar to our previous research [ 18] .The white grain is determined to be a brittle intermetallic of AlNb2 from the identification by EBSD,which is harmful to the mechanical property of the resultant joint.Figure 4b shows another typical micros true ture of TiAl/Ti2AlNb bonded interface with pure Ti interlayer at the same condition.Distinctly,there is no brittle intermetallic formed at the bonding interface,and the DLs are more uniform compared to those in Fig.4a.Moreover,Table 2 shows the quantitative analysis results of the points in Fig.4b determined by EMPA.It is also shown that there are no intermetallic compounds produced at the bonding interface.In particular,Point 9 has an extra Mo content of?2.55 at%,which is labelled as p/B2phase close to TiAl alloy.

Fig.3 Schematic diagram of EBSD and mechanical performance test:a EBSD test coordinate system and b tensile specimen

Fig.4 Typical SEM images of TiAl/Ti2AlNb bonded interface at 1000℃-20 MPa-90 min:a without an interlayer and b with pure Ti interlayer

Moreover,two layers are produced at the bonding interface from SEM image.One is a single gray strip layeradjacent TiAl alloy,and the other is a layer containing white band inlaid with gray needles and black equiaxed phase near the Ti₂AlNb alloy.Figure 5 shows the line and plane scan of the major element across the bonding interface at 1000℃-20 MPa-90 min.There is an interdiffusion of major elements obviously across the bonding interface,as shown in Fig.5a.Figure 5b shows the plane distribution of the major elements within the white rectangle.There is an obvious enrichment of Nb element at the bonding interface,which is the main reason for the formation of"beach"micros true ture at the interface between the two DLs.In addition,the thickness of the DLs is approximately twice as thick as the joint without an interlayer approximately,as shown in Fig.4.This will be discussed in Sect.3.3.

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Table 2 Elemental quantitative analysis results of points in Fig.4b determined by EMPA

To further confirm the products within DLs,EBSDtechnology was implemented to research the crystallographic characterization of the interfacial microstructure.The phase distribution of the bonded joint is shown in Fig.6a,where theγphase is presented as blue,βphase as yellow,α₂ phase as red and O phase as green,respectively.Also,the low angle grain boundaries (LAGBs) were measured with a misorientation (θ) of grain boundaries between 2^°and 15^°(light blue line),and theθvalue of high angle grain boundaries (HAGBs) is higher than 15^°(black line).The boundaries with 57^°<θ<63^°are considered as twin boundaries (TB)(purple line).One can see that both gray acicular phase and gray equiaxed phase (Points 3and 4) areα₂ phases.There is almost no LAGB at the bonding interface.It proves that dynamic recrystallization(DRX) occurs on the interfacialα₂ phase ultimately.It is conducive to dislocation movement and the evolution of the contact interface,and the density of dislocation at the DRX region is reduced as the corresponding GND map shows in Fig.6b.The similar phenomenon has also occurred in the bonding of Ti alloy [ 27] and TiAl alloy [ 28] .

Fig.5 EPMA analysis of major elements across bonding interface at 1000℃-20 MPa-90 min:a line scan of black line and b plane distribution within white rectangle

Furthermore,the grain size of the interfacialα2 phase is perse.Scilicet,the grain size of theα2 phase in the middle (I-α2) of the bonding interface is the smallest,on account of being induced by DC magnetron sputtering.The grain size of theα2 phase adjacent to Ti2AlNb alloy(Ti2AlNb-α2) is larger than that of the intermediateα2phase.The grain size of theα2 phase adjacent to TiAl alloy(TiAl-α2) is the largest.The I-α2 is equiaxed grain,while the Ti2AlNb-α2 and TiAl-α2 are elongated grains (Fig.6a).It can be revealed that the crystal nucleates at the I-α2 and grows toward the two based alloys with the process of the bonding.This behavior is the criterion to measure the thickness of the DLs in EBSD maps,following which the thickness of every reaction layer was measured.The diffusion direction of the interlayer Ti atom is reciprocal to the two based alloy at the initial stage and tend to be gentle gradually,as shown in Fig.7c.In particular,the phase transformation is found at the side of Ti2AlNb based alloy,and the newα2 phase is always close to the O phase,asshown by black rectangle region (A and B) in Fig.6a.This phenomenon indicates that lots of O phases form thisα2phase.The orientation relationship between equiaxed/acicularα2 phase and O phase can be described by [ 1] 。//, [ 17] ,labelled in Areas A and B,separately,as shown in Fig.6a.And Fig.6c,d shows the corresponding pole figures of two phases,indicating that the phase transformation from O (rim O and lamella O) toα2 has occurred.This phase transformation behavior will promote the microstructural evolution of the bonding interface.

Figure 7 shows TEM bright-filed image of the hightemperature TiAl/Ti2AlNb bonding interface,in which the black line shows the recrystallized grain of I-α2 interlayer.To further study the I-α2 interlayer,the magnification of region 1 in Fig.7a was conducted.Figure 7b shows the corresponding HRTEM image,there is no other nanophase in I-α2 interlayer based on fast Fourier transformation(FFT) patterns of Regions A,B and C.

Fig.6 EBSD maps of bonding interface at 1000℃-20 MPa-90 min:a phases map,b GND map,and c,d pole figures

Fig.7 TEM analysis of bonding interface:a TEM bright-field micrograph and b HRTEM image of Region 1 (insets being FFT images);c line scan of green area in a

Besides,to determine the direction of interface growth during the diffusion bonding,the grain orientation distribution of TiAl/Ti₂AlNb joint is shown in Fig.8.Figure 8a-c shows the grain orientation distribution of normal direction (ND),bonding direction (BD) and transverse direction (TD) at the bonding interface,respectively.According to Fig.8,it can be found out that there exists hardly preferred orientation of the grains across the binding interface along with ND,BD or TD directions.Moreover,the full pole figures (Fig.9a) and inverse pole figures (Fig.9b) ofα₂ grains in the bonding interface were investigated,which shows that the texture of interfacialα₂grains is free.The grain grows freely during the bonding due to the effect of nanointerlayer,i.e.,there are a number of grain boundaries at the bonding interface to furnish the vacancy and promote the diffusion path for atoms.In addition,when the grain boundary is parallel to the stress,the stress direction is identical to the BD during the diffusion bonding,and the formation energy of the vacancy is high.Nevertheless,the formation energy of the vacancy is lower as the grain boundary is perpendicular to stress.Hence,the vacancy will assemble along the BD.Consequently,both bar-like grains and sawtooth phase interface of Ti₂AlNb-α₂ and TiAl-α₂ are formed at the bonding interface,respectively.It can be inferred that the growth of the bonding interface is controlled by vacancy diffusion.

3.2 Microstructural evolution of bonding interface

Figure 10 reveals the impact of bonding time on the interfacial microstructure of TiAl/Ti₂AlNb joints bonded using a pure Ti interlayer at 1000℃under a pressure of20 MPa.It is observed that the same interfacial microstructure is formed in the joints and the thickness ofthe DLs gradually increases with the increase of bonding time,as shown in Fig.10a-c.Besides,both the thicknesses of the reaction layer adjacent Ti₂AlNb alloy and the reaction layer adj acent TiAl alloy have the same trend with the total thickness.With a lower bonding time of 60 min,as shown in Fig.10a,some coarse white clusters together with many black refined particles are stacked on the clad interface,which shows an insufficient diffusion.Further,with the increase of the bonding time to 90 min,Fig.10b shows that the cohesive interface is smooth,the black particles are dispersed gradually,and the interfacial microstructure is uniform.As shown in Fig.10c,the bonding interface is almost straight when the bonding time is 120 min.However,it is found that the grain of DLs adjacent TiAl alloy is coarse,which is detrimental to the mechanical properties.

3.3 Grow kinetics of diffusion reaction layers

For the sake of clarifying the growth behavior of the DLs detailly,the color-coded phase maps of the bonding joints at1000℃-20 MPa with different bonding time obtained from EBSD analysis were conducted,as shown in Fig.11.As mentioned above,the DLs can be pided into two layers from SEM,but the DLs should be broken into four layers from EBSD analysis in further research.Here,the thickness of DLs was measured by the minimum of diffusion distance according to the phase boundary,which is indicated by the black line in Fig.11.According to Fig.11,the DLs contains four layers,including (Ⅰ) single coarseα2phase adjacent TiAl alloy,(Ⅱ) single refined phase at the bonding interface,(Ⅲ) equiaxed/acicularα2 phase embedded inβphase adjacent Ti2AlNb alloy and (Ⅳ) both equiaxedα2 phase and acicular O phase embedded inβphase adjacent Ti2AlNb alloy,respectively.Of course,there has an error of artificial microscale measurement.The diffusion thicknesses of four layers are marked in Fig.11,which are averaged by three measures.

Fig.8 Grain orientation distributions of TiAl/Ti2AlNb joint:a ND,b BD,and c TD

Fig.9 a Full and b inverse pole figures ofα2 grains in DLs

Fig.10 SEM images of bonded joint at 1000℃-20 MPa with different bonding time:a 60 min,b 90 min,and c 120 min

Fig.11 EBSD color-coded phase maps of bonding joints at 1000℃-20 MPa with different bonding time obtained:a 60 min,b 90 min,and c 120 min (color-coded agreeing with former)

Figure 11a indicates the thickness of DLs at different bonding time.Note that the thickness of each layer of thejoints increases as the bonding time increases,as also shown in Fig.12a.In order to illuminate the layer growth behavior,it is crucial to analyze the growth kinetics of the four reaction layers.The thickness of a reaction layer by elemental diffusion at a given temperature can be expressed by the following relationship [ 29] :

Fig.12 a Evolution of diffusion thicknesses of interfacial reaction layers at different bonding time;b linear regressions according to Eq.(2);c linear regressions according to Eq.(3)

where X is the thickness of reaction layer,k is the reaction rate constant,t is diffusion time,and n is reaction kinetic exponent.Distinct value of n represents the different growth mechanism of reaction layers,standard diffusion limit growth is expressed when n is 0.5,and n>0.5suggests an interface-controlled growth [ 30] .The DLs were formed by the diffusion of atoms as the above discussion.Therefore,Eq.(1) can be written by:

The experimental data shown in Fig.11 are processed into Eq.(2) for linear regression analysis,and the result is shown in Fig.12b.As shown in Fig.12b,the reaction kinetic exponent (n) for different layers are 0.46,0.56,0.13and 0.31,respectively.As for LayersⅠandⅡ,the value of n is close to 0.5,showing that the layers are controlled by a standard diffusion.However,for LayersⅢandⅣ,the value of n is lower than 0.5,which may be caused by the phase transformation at the based Ti₂AlNb alloy.Also,it is necessary to take the experimental error into account.As a result,it is inferred that the layer growth belongs to a standard diffusion limited growth.

Based on Eq.(3),the experimental data could be expressed as a linear analysis,as shown in Fig.12c.Hence,it can be given that the values of k for RegionsⅠ,Ⅱ,ⅢandⅣare 1.22×10-6,1.27×10-6,2.60×10-7 and7.70×10-7 m·s-1/2,respectively.

3.4 Mechanical properties of bonding joints

The RT tensile strength of the TiAl/Ti/Ti₂AlNb DB joint obtained at different bonding time is shown in Fig.13.It is clear that the tensile strength increases firstly and then decreases with the increase of bonding time.Figure 13reveals that the RT tensile strength is relatively lowbecause of an insufficient diffusion with a short bonding time of 60 min.There is a visible improvement of tensile strength with the bonding time increasing to 90 min.And the maximum tensile strength of 281 MPa is achieved,which is about twice as strong as that of the joint without any interlayer marked by the red line in Fig.13.Table 3shows the comparison of the strength of the dissimilar diffusion bonding joint of TiAl or Ti₂AlNb alloy with/without the interlayer.Note that the addition of the interlayer can enhance the strength of the joint as the results in Refs. [ 19, 21, 22, 23] .However,the strength is also needed to be strengthen.Apart from this,the composite interlayer is beneficial for the weld of the dissimilar materials,which has a positive effect on interfacial microstructure control and the performance of bonding joint [ 31] .Increasing the bonding time leads to a decrease of tensile strength due to the growth of grains.Also,all the samples were fractured at the bonding interface,which may be caused by the hard and brittleα₂ layer in the reaction layers.In conclusion,TiAl alloy and Ti₂AlNb alloy can be bonded at 1000℃-20 MPa-90 min by adding a nano Ti interlayer and the RTtensile strength improves obviously.Furthermore,the design of composite interlayer is another important direction for the bonding of TiAl and Ti₂AlNb alloy.The studies of the design of composite interlayer to further improve the properties of the dissimilar bonding joint will be focused in the future.

Fig.13 Tensile strength of TiAl/Ti2AlNb DB joint using pure Ti interlayer

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Table 3 Strength of dissimilar diffusion bonding joint of TiAl or Ti2AlNb alloy

Tensile strength was marked by T,and S represented shear strength

Fig.14 Schematic diagrams of microstructural evolution of TiAl/Ti2AlNb DB joint:a before DB,b under DB,and c after DB

3.5 Diffusion bonding mechanism

To study the diffusion bonding of the TiA alloy and Ti2AlNb alloy,the pure Ti layer fabricated by DC magnetron sputtering is used.The pure Ti interlayer has a significant effect on micros true tural evolution and mechanical performance.The pure Ti interlayer could buffer the mechanical stress at the faying interface before diffusion bonding,thus decreasing the internal stress in the bonding joint.During the bonding process,many grain boundaries provided by the nano pure Ti interlayer can promote the diffusion of atoms and voids sealed,which is importance for the microstructural evolution of the DLs.Therefore,the resultant joint will be produced after furnace cooling.The diffusion bonding mechanism is schematically illustrated in Fig.14.Figure 14a shows the state before diffusion bonding.And the interlayer has a great number of grain boundaries to adjust the deformation and facilitate diffusion of the element,as shown in Fig.14b,during the DB processing.Besides,a metallurgical bonded joint with excellent properties is formed in Fig.14c.

4 Conclusion

TiAl alloy and Ti2AlNb alloy were successfully bonded using a pure Ti interlayer by diffusion bonding.The results show that the DLs mainly consisted of four characteristic layers:(Ⅰ) single coarseα2 phase adjacent TiAl alloy,(Ⅱ)single refinedα2 phase at the bonding interface,(Ⅲ)equiaxed/acicularα2 phase embedded inβphase adjacent Ti2AlNb alloy and (Ⅳ) both equiaxedα2 phase and acicular O phase embedded inβphase adjacent Ti2AlNb alloy.

The thickness of the four layers increased with the increase of the bonding time.The growth of DLs is limited by diffusion growth and the reaction rate constant (k) for RegionsⅠ,Ⅱ,ⅢandⅣare 1.22×10-6,1.27×10-6,2.60×10-7 and 7.70×10-7 m.s-1/2,respectively.And the orientation of the interfacialα2 phase grows up without texture.The RT tensile strength of the joints increased initially,and after that,it decreased with the bonding time increasing.The RT tensile strength of 156 MPa is obtained without the interlayer,and RT tensile strength increases to281 MPa with a nano interlayer at 1000℃-20 MPa-90 min.

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