Materials and Design 52 (2013) 359–366

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Materials and Design

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Impulse pressuring diffusion bonding of a copper alloy to a stainlesssteel with/without a pure nickel interlayer

0261-3069/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.matdes.2013.05.057

⇑ Corresponding author. Tel./fax: +86 23 65127306.E-mail address: xinjianyuan@cqu.edu.cn (X. Yuan).

Xinjian Yuan ⇑, Kunlun Tang, Yongqiang Deng, Jun Luo, Guangmin ShengCollege of Materials Science and Engineering, Chongqing University, No. 174, Shazheng Street, Shapingba District, Chongqing 400044, PR China

a r t i c l e i n f o

Article history:Received 18 December 2012Accepted 17 May 2013Available online 31 May 2013

Keywords:Impulse pressuringDiffusion bondingCopper alloyStainless steelMechanical property

a b s t r a c t

Impulse pressuring diffusion bonding of a copper alloy to a stainless steel was performed in vacuum.Using Ni interlayer of 12.5 lm, the joint produced at 825 �C under 5–20 MPa for 20 min exhibited lowerstrength, which could result from the insufficient thermal excitation and plastic deformation. At 850 �Cunder 5–20 MPa for 5–20 min, the strength of the joint improved with time. An optimized joint strengthreached up to 217.2 MPa. Fracture occurred along the Cu–Ni reaction layer and the Ni layer and almostplastic fracture was confirmed by extensive dimples on the fracture surface. Using the interlayer of50 lm, the fracture surface was similar. Without Ni assistance, under the same bonding condition, thejoint strength was about 174.2 MPa. The lowered strength might be attributed to the appearance of someunbonded zones in the joint. Lots of brittle fracture areas appeared on the fracture surface.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Because of the beneficial features of the bimetal, the demand forjoints of Cu alloy/stainless steel has increased in industrial applica-tions from nuclear, automobile, rail and aviation industries tosmaller, more commonly used products [1]. For instance, in thepower-generation plants, the Cu alloy/stainless steel combinationshave been widely used and primarily employed for the in-vesselcomponents of international thermonuclear experimental reactor(ITER) [2–5]. ITER has been constructed to produce efficient andcommercial energy from fusion by safe operation of a reactor.Based on the successful applications of austenitic stainless steelin the nuclear environments of light water and faster breeder reac-tors, it is selected as a primary structural material for first wall/blanket system. In view of high conductivity and thermal stressresistance of Cu alloys, they are selected as a heat sink to meetthe particular demand that heat fluxes generated in the divertorshould dissipate uniformly [5]. Joining of Cu alloy to stainless steelhas become a challenging task due to their obvious mismatches,such as thermomechanical properties, physical properties andchemical properties.

Such conventional welding processes as shielded metal arcwelding, gas tungsten arc welding, gas metal arc welding and sub-merged arc welding make Cu alloy and stainless steel difficult toobtain defect-free dissimilar joints [6]. To improve the microstruc-ture and properties of Cu alloy/stainless steel joint, the bimetal hasbeen produced by electron-beam welding [7,8], laser welding [6,9],

friction welding [10,11], explosive welding [3,4,12,13] and diffu-sion bonding [1,2,14–17]. Among these joining methods, diffusionbonding is an attractive technique for forming this bimetal whileminimizing deleterious effects (such as crack, distortion and segre-gation) on material properties [3,14,15]. Diffusion bonding is a so-lid-state joining process, which is conducted under an exertingpressure at 0.7–0.9 of the absolute melting temperature of thematerials to minimize macroscopic deformation and allows con-tacting surfaces to be joined [18]. Generally, during any direct dif-fusion bonding of dissimilar materials, poor mechanical propertiesmight result from the residual stress produced by the related mis-match in the material linear expansion coefficients and the formedmicro-voids and micro-cracks owing to the difference in intrinsicdiffusion coefficient of alloying species. For example, Yilmaz andAksoy have found that micro-voids and micro-cracks are formedin the bonding region and cannot vanish even though the joiningtime is increased to 30 min. Furthermore, these defects weakenthe mechanical and physical properties of the joint [16].

To reduce the negative effects, increase the bonding efficiencyand improve the joint properties, two efforts were made in this pa-per. On the one hand, impulse pressuring diffusion bonding wasapplied as the joining process. Impulse pressuring diffusion bond-ing process, during which an impulse pressure substitutes for aconstant pressure, has been employed successfully to join dissim-ilar materials [19,20]. For instance, successful studies have been re-cently reported on impulse pressuring diffusion bonding oftitanium alloy to stainless steel. The tensile strength of the jointhas been enhanced to 321 MPa and the effective bonding timehas been decreased to 180 s [19]. However, little work has beendone regarding impulse pressuring diffusion bonding in copper

Fig. 1. Schematic diagram of impulse pressuring diffusion bonding process.

Fig. 2. Schematic drawing for tensile testing.

360 X. Yuan et al. / Materials and Design 52 (2013) 359–366

alloy-stainless steel bimetal. On the other hand, nickel was se-lected as an intermediate. That is because of good mutual solid sol-ubility between nickel and iron as well as between nickel andcopper [17]. Furthermore, the corrosion resistance of the bondmade using nickel as an intermediate is better as compared tothe joints obtained using other interlayer [14].

Therefore, the objective of the present research is to investigatemicrostructural characteristics and mechanical properties duringimpulse pressuring diffusion bonding of a copper alloy to a

Fig. 3. SEM micrographs and EDS line scan results (the black lines showing the line scan5 min (a), 10 min (b) and 20 min (c).

stainless steel with a pure nickel foil interlayer. Impulse pressuringdiffusion bonding without nickel assistance was carried out as areference to do comparative analyses.

2. Experimental procedure

304L stainless steel and pure copper (99.9 wt.%) cylinders, mea-suring U12 mm � 35 mm, were machined from as-received rodsand used as the base materials. The nominal chemical compositionof the stainless steel in weight percent was 0.02% C, 1.26% Mn,19.35% Cr, 0.39% Mo, 8.15% Ni, 0.42% Si, 0.20% Cu, 0.024% P,0.004% S and the balance was Fe. Pure Ni (99.99 wt.%) foils with12.5 lm and 50 lm were used as the interlayer.

The surfaces of the sectioned coupons were ground on SiC paperto a 2000-grit finish and then polished by a 1 lm polycrystallinediamond suspension to ensure an adequate mating surface. Priorto impulse pressuring diffusion bonding, the specimens to bebonded together were decreased with acetone in an ultrasonicbath for 60 min.

The interlayer foil was sandwiched between the mating sur-faces of the base metals. A series of bonding tests were performedin a Gleeble-1500D under a vacuum of approximately 5 � 10�2 Pa.A schematic diagram of impulse pressuring diffusion bonding pro-cess is shown in Fig. 1. Based on the pre-experiments and the Ref.[19], a uniaxial impulse load of minimum pressure (5 MPa) andmaximum pressure (20 MPa) was applied along the longitudinaldirection of the assemblies, and the frequency of impulsion was0.5 Hz. The number of impulsion corresponds to the bonding time.During bonding, both the heating rate and the cooling rate were5 �C/s. After bonding operation, the bonded specimens were an-nealed for 1 h at 400 �C in vacuum furnace to eliminate joiningstress and produce optimal structure. Specimens for tensile testingwere machined using a standard tensile test in accordance to

trace) for typical cross-sections prepared with Ni interlayer of 12.5 lm at 850 �C for

Fig. 4. Average width of interdiffusion layers changed as a function of bondingtime.

Fig. 6. SEM micrograph (a) and EDS line scan results (b) (the dark dot line showingthe line scan trace) for typical cross-section obtained using Ni interlayer in athickness of 50 lm under 5–20 MPa impulse pressure at 850 �C for 20 min.

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ASTM: E8/E8M-11. The tensile testing is shown schematically inFig. 2. And then the machined specimens were evaluated on an In-stron-1342 mechanical testing machine with a constant strain rateof 0.5 mm/min.

For microstructural examinations, specimens were sectionedfrom the diffusion-bonded joints. Ground and polished cross sec-tions were etched for 5–10 s in aqua regia (60% HCl + 20%HNO3 + 20% H2O). Copper alloy base materials were etched forabout 60 s in a mixture of 5 g FeCl3, 50 ml HCl and 100 ml distilledwater (H2O). Microstructural observations were conducted usingan optical microscope (OM) and a NOVA-400 field emission scan-ning electron microscope (FE-SEM). The chemical compositionwas analyzed by means of energy-dispersive X-ray spectroscopy(EDS).

3. Results and discussion

3.1. Cross-sectional microstructure of joint

Fig. 3 illustrates typical cross-sectional SEM images and EDSline scan results of Cu/304L joints prepared using impulse pressur-ing diffusion bonding with Ni interlayer in a thickness of 12.5 lmunder 5–20 MPa impulse pressure at 850 �C for 5, 10 and 20 min.From Fig. 3, four conclusions can be drawn.

Firstly, for three bonding conditions, there were gradual in-creases in the Ni content and gradual decreases in the Fe, Cr andCu contents at the 304L-Ni and Ni–Cu interfaces. At these inter-faces, no composition platforms were observed, which conveysthat intermetallic compounds were not formed.

Secondly, the remainder of Ni interlayer was found in the jointcenter area. This means that the existence of Ni prevented theinterdiffusion of Fe, Cr and Cu.

Fig. 5. OM images of the copper alloy exposed for

Thirdly, in the case of 5 min, some micro voids in an irregularshape could be seen clearly in the interface zone. This may beattributed to the fact that the diffusion coefficients of Fe, Cr, Niand Cu are different and then a flux imbalance across the interfacesis originated. Another reason might be that the original porositiescannot be eliminated by the contact of surfaces and the deforma-tion and flow of metals (especially Ni foil) due to lower holdingtime. Mechanical properties of joint are adversely affected by thesegenerated weak void zones [4,14]. Similar occurrences of Kirken-

20 min at 825 �C (a), 850 �C (b) and 875 �C (c).

Fig. 7. SEM micrograph (a) and EDS result (b) (the white dot line showing the scantrace) for typical cross-section made without Ni interlayer at 850 �C for 20 min.

Fig. 8. Variations in the tensile strength of the joints prepared as functions ofbonding time (a), bonding temperature (b) and interlayer thickness (c) (0 lmthickness interlayer corresponding to no interlayer).

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dall void were found within the diffusion-bonded interfaces[1,4,14,15]. By comparison with other studies, the amount of microvoids is very small in the present research. This indicates that dif-fusion bonding with impulse pressure is beneficial to the reductionof micro voids. Furthermore, it could be observed from Fig. 3 thatthese micro voids lessened as a function of time, since the defor-mation of metals and the contact of surfaces increase with time.As a result, a sound joint could be obtained.

Finally, the interdiffusions of Cu and Ni elements, Fe, Cr and Nielements can be demonstrated by their concentration profilesacross the interface zones of the joints. The average thickness ofinterdiffusion layer in the interface areas was measured usingSEM micrographs and plotted against the bonding time, as givenin Fig. 4. From Fig. 4, the thickness of reaction layer widened withthe increase of time. It is perhaps worthy of note that the interdif-fusion depth between Ni and Cu at the Ni–Cu interface was largerthan that between Fe, Cr and Ni at the 304L-Ni interface. Thismight result from the diffusion of Ni and Cu is faster as comparedto the diffusion of Fe and Cr. For instance, the diffusion coefficientof Cu (5 � 10�14 m2 s�1 at 900 �C) is higher than that of c-Fe(3 � 10�17 m2 s�1 at 900 �C) according to Refs. [1,14,21].

To understand the effect of temperature on the microstructureof the joint and the copper alloy base metal, impulse pressuringdiffusion bonding of 304L to Cu alloy was performed as a functionof temperature from 825 �C to 875 �C using Ni interlayer of12.5 lm under 5–20 MPa impulse pressure for 20 min. The resultswere similar to the cases that mentioned with time. There werefew micro-voids in the bonding region. As the bonding tempera-ture was increased, the interdiffusion layer broadened. Fig. 5 shows

OM images of the copper alloy base material exposed at 825 �C,850 �C and 875 �C for 20 min. Exposure at high temperaturescoarsened the microstructure of the copper alloy. That is whypost-bond heat treatment was carried out. Re-heat treatment ofthe specimens can made optimal structures and relieve bondingstress.

Besides, impulse pressuring diffusion bonding of 304L to Cu al-loy was also conducted using interlayer with 50 lm. Fig. 6represents typical cross-sectional SEM images and EDS line scanresults of Cu/304L joints obtained using Ni interlayer in a thicknessof 50 lm under 5–20 MPa impulse pressure at 850 �C for 20 min.

Table 1Tensile strength of the copper alloy subjected to high temperature exposure for20 min at 825 �C, 850 �C and 875 �C.

Exposure temperature (�C) Tensile strength (MPa)

825 241.5850 238.8875 233.1

Table 2Chemical compositions (analyzed by EDS) of the micro-zones in Fig. 9.

Micro-zones Elemental composition (at.%)

Fe Cr Ni Cu

Zone 1 50.15 11.00 38.85 –Zone 2 68.75 9.96 21.29 –Zone 3 3.17 – 96.83 –Zone 4 – – 14.78 85.22Zone 5 – – 100.00 –Zone 6 – – 15.81 84.19Zone 7 – – 100.00 –Zone 8 – – 23.90 76.10

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Diffusion-bonded joint was also produced. The width of the resid-ual Ni interlayer in the bonding region was relatively larger.

To make a comparison, impulse pressuring diffusion bonding of304L to Cu alloy without interlayer was also inspected. Fig. 7 pre-sents typical cross-sectional SEM images and EDS line scan resultsof Cu/304L joints made using impulse pressuring diffusion bondingwithout Ni under 5–20 MPa impulse pressure at 850 �C for 20 min.From Fig. 7, due to the interdiffusion of Fe, Cr and Cu elements, athin transition zone including Fe, Cr and Cu with a thickness ofabout 3 lm was seen in the interface. Additionally, there weresome unbonded zones in the interface area. These unbonded zonesmight come from the original porosities and Kirkendall voids. Dueto different diffusion coefficients, Micro-voids were also observedin the Cu-304 stainless steel couple produced by direct diffusionbonding without interlayer [1,16].

3.2. Tensile strength of joint

Variations in the tensile strength of the joints prepared as func-tions of bonding time, bonding temperature and interlayer thick-ness were summarized in Fig. 8. It can be seen from Fig. 8a thatan increase of the bonding time caused an enhancement of thejoint strength. For a short time, the tensile strength of the jointwas the lowest. That is because the reaction layer is thin and thereare some voids present in the interface area. As the bonding timewas increased to 10 min, the reaction layer widened and the voidsshrank, and thus the joint tensile strength improved to 196.9 MPa.With further increase in the bonding time to 20 min, a sound jointwas made, and consequently the tensile strength was up to217.2 MPa, which is approximately 79% of the Cu substratestrength (275 MPa) and 91% of the tensile strength of the Cu alloysubjected to high temperature exposure at 850 �C, as given in Table1. From Table 1, after exposure for 20 min at 825 �C, 850 �C and875 �C, the tensile strength of the copper alloy base metal was241.5 MPa, 238.8 MPa and 233.1 MPa, respectively. As describedin Ref. [22], tensile property of copper alloy reduced gradually withtemperature when the treating temperatures were above the

Fig. 9. Typical fracture surface morphologies from the bonds conducted with N

recrystallization temperature of copper alloy. This indicates thathigh temperature exposure causes a decrease in the tensilestrength, which makes the bond strength difficult to furtherimprove.

Fig. 8b describes the effect of bonding temperature on the ten-sile strength when the other parameters were kept constant. Whenthe bonding temperature was in the range of 825–850 �C, strengthrose as the temperature increased. At low bonding temperature,the thermal excitation of the mating surfaces might be not compa-rably sufficient [14]. With increasing temperature, the plasticdeformation of the materials bonded could enhance and the atomicdiffusivity could quicken, and then the joint strength improves. At850 �C, the tensile strength of the bond reached to a maximum va-lue. With further increase in the temperature to 875 �C, the bondstrength reduced slightly, which may be attributed to the graingrowth and the microstructure softening at a higher temperature.

From Fig. 8c, the strength of the joint exhibited a relativelysmaller value, when the interlayer thickness was 50 lm. Whenthe thickness of the interlayer enlarges, the plastic deformationof the interlayer metal could become insufficient. This should beresponsible for lower strength of joint produced using the 50 lmthickness interlayer.

In the case of impulse pressuring diffusion bonding without Nifoil interlayer for 20 min, the tensile strength of the joint was about174.2 MPa, which is obviously less than the joint strength obtainedwith Ni foil at the same bonding condition. The appearance of someunbonded zones may be responsible for the smaller value for im-pulse pressuring diffusion bonding without Ni foil interlayer.

The optimum value obtained in this study is larger than thejoint strength of stainless steel to Cu alloy joined using explosivewelding and hot isostatically pressed bonding [23]. Additionally,Sabetghadam et al. have examined conventional diffusion bondingof a 410 stainless steel to a pure Cu using a nickel interlayer and it

i interlayer of 12.5 lm at 850 �C for 5 min (a), 10 min (b) and 20 min (c).

Fig. 10. Typical fracture surface morphologies from the bonds carried out with Ni interlayer of 12.5 lm for 20 min at 825 �C (a) and 875 �C (b).

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has been reported that the bonding temperature and the holdingtime for obtaining the maximum value of shear strength(145 MPa) are 900 �C and 60 min, respectively [14]. By contrast,the bonding temperature lowers by 50 �C and the bonding timeshortens to only one third of it in this paper. The achievementsof high efficiency and strength are possible due to impulse pressur-ing diffusion bonding method and Ni interlayer assistance.

3.3. Fracture characteristics

Fig. 9 shows fracture surface morphologies of the specimensprepared using impulse pressuring diffusion bonding with Ni foilof 12.5 lm under 5–20 MPa impulse pressure at 850 �C for 5, 10and 20 min. EDS analysis results are given in Table 2. Fig. 10 de-scribes fracture surface morphologies of the specimens producedusing impulse pressuring diffusion bonding with Ni foil of

Fig. 11. Typical fracture surface morphologies from the bonds produced at 850 �Cfor 20 min with Ni interlayer of 50 lm.

12.5 lm under 5–20 MPa impulse pressure for 20 min at 825 �Cand 875 �C. Fig. 11 represents fracture surface morphologies ofthe specimens obtained using impulse pressuring diffusion bond-ing with Ni foil of 50 lm under 5–20 MPa impulse pressure at850 �C for 20 min. Fig. 12 exhibits fracture surface morphology ofthe sample made using impulse pressuring diffusion bonding with-out Ni foil under 5–20 MPa impulse pressure at 850 �C for 20 min.EDS analysis results are given in Table 3.

Fig. 12. Typical fracture surface morphology from the joint performed without Nifoil at 850 �C for 20 min.

Table 3Chemical compositions (analyzed by EDS) of the micro-zones in Fig. 12.

Micro-zones Elemental composition (at.%)

Fe Cr Ni Cu

Zone 1 30.02 8.91 3.42 57.65Zone 2 4.59 – – 95.41

Fig. 13. Stress–strain curves upon loading during tensile testing for joints prepared with Ni interlayer of 12.5 lm at 850 �C for 10 min (a) and 20 min (b).

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From Fig. 9a, in the case of 5 min, failure took place along twodifferent regions, which are termed the black and gray region (zone1 and zone 2) and the dimple region (zone 4). According to Table 1,the black and gray region is Fe–Ni solid solution. SEM/EDS analysissuggests that the gray solid solution (zone 1) is enriched in Ni,while the black solid solution (zone 2) is relatively depleted ofNi. This indicates that zone 1 is adjacent to the Ni interlayer andzone 2 is near to the stainless steel. Zone 4 could be regarded asCu–Ni solid solution based on Table 1 and this zone with a patternof dimple reveals the fracture was plastic fracture. By comparisonwith the black and gray region, the area of the dimple region on thefracture surface was small. Because the low bonding temperaturecan cause a minimum thermal activation and the minimized extentof solute diffusion [14], the amount of Cu–Ni solid solution formedis limited. In this condition, fracture occurred along the stainlesssteel–Ni foil interface and the Cu alloy–Ni foil interface. But, theproportion of plastic fracture was lower significantly.

With a subsequent increase of holding time to 10 min, lots ofdimples (zone 6) and a few gray zones (zone 5) were observedon the fracture surface. On the one hand, the area of zone 6 notablyincreased by prolonging time and this zone is also Cu–Ni solidsolution. On the other hand, Fe–Ni solid solution for 5 min was re-placed by pure Ni gray zone. The area of zone 5 on the fracture sur-face was markedly smaller than that of zone 6. Zone 6 with adimple pattern exhibited plastic fracture. By contrast, non-plasticzone 5 may be considered as a relatively weak zone.

As the bonding time was prolonged to 20 min (Fig. 9c), threeconclusions could be made as compared to 10 min. First, the per-centage of the relatively weak zone decreased on the fracture sur-face. Second, almost plastic fracture was found after tensile testing.Third, the concentration of solute in Cu–Ni solid solution increasedaccording to Table 2.

Based on the analyses of fracture surfaces, the lowest tensilestrength of the joint results from the small proportion of plasticfracture occurred for 5 min. For 10 min, an increase in the area ofdimple zone causes an improvement in the tensile strength. But,the strength value for 10 min is less than that for 20 min. Thismight be explained in the following two reasons. First, becauseof the presence of the relatively weak zones, crack initiation is in-duced preferentially in these zones. This could be confirmed by theresults of Fig. 13. There are some undulate segments in the stress–strain curve. These undulate segments demonstrate that cracksformed upon loading. The undulate segment corresponds to therelatively weak micro-zone. The amount of the relatively weak mi-cro-zone for 10 min (Fig. 13a) was notably larger than that for20 min (Fig. 13b). Second, an increase in the concentration of sol-

ute in Cu–Ni solid solution might enhance the weldability byincreasing 10 min to 20 min.

From Fig. 10a, the chief character of the fracture surface wasparallel striation. Some shallower gray zones exhibited dimple pat-tern. A reduction in the tensile strength at 825 �C may result fromthe presence of this stripe-shaped morphology having a strongdirectionality. In the case of 875 �C (Fig. 10b), the fracture surfacewas similar to the case of 850 �C and had lots of dimples after ten-sile testing. When using the foil of 50 lm as the interlayer, theappearance of a lot of dimples conveys that the fail mode wasmainly plastic fracture mode and shows marked similarities withFig. 9c.

In the case of bonding without Ni foil interlayer, it can be seefrom Fig. 12 that there was different morphology on the fracturesurface. Fail occurred along the Cu–Fe solid solution and the Cu al-loy. The fractography of the joint conveys failure mode involved aductile pattern and a brittle pattern. A great amount of brittle frac-ture area appeared on the fracture surface. Thus, the tensilestrength of the joint was comparatively low. The fracture surfacealso included ductile fracture area and brittle fracture area andfracture took place with micro-cracks propagated from voids dur-ing diffusion bonding of Cu alloy and stainless steel without Niinterlayer [1,16].

4. Conclusions

In the present study, the effects of bonding parameters on themicrostructure and mechanical properties of the joints producedby impulse pressuring diffusion bonding of a copper alloy to astainless steel were investigated. From this research, the followingconclusions can be drawn:

(1) The presence of Ni, acting as a transition and barrier layer,impedes the interdiffusion of Fe, Cr and Cu elements. With-out Ni foil, a thin zone containing Cu, Fe and Cr with a thick-ness of about 3 lm was found in the interface area.

(2) Diffusion bonding with impulse pressure benefited lesseningmicro-voids and their amount. With Ni interlayer of 12.5 lmat 850 �C under 5–20 MPa impulse pressure, micro-voidsdecreased obviously and thus the tensile strength of thejoint increased as function of bonding time. When the timewas prolonged to 20 min, the joint strength reached up to217.2 MPa. Fracture took place along the Cu–Ni reactionlayer and the remnant Ni layer. Additionally, a great numberof dimples convey that almost plastic fracture was foundafter tensile testing.

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(3) With Ni interlayer of 12.5 lm under 5–20 MPa impulse pres-sure for 20 min, the joint strength experienced a hill-shapedcurve change from 825 �C to 875 �C. At a lower temperature,the thermal excitation of the mating surfaces and the plasticdeformation of the materials may be not comparably suffi-cient. At a higher temperature, the grain growth and themicrostructure softening might result in the slight reductionof the bond strength.

(4) Using Ni interlayer of 50 lm at 850 �C under 5–20 MPaimpulse pressure for 20 min, lower strength of the jointshould be attributed to the insufficient plastic deformationof the interlayer metal due to larger thickness of theinterlayer.

(5) Without Ni foil, under the same bonding condition, the exis-tence of some unbonded zones in the joint area brought outrelatively lower tensile strength (174.2 MPa). The fractogra-phy indicates that failure mode included a ductile patternand a brittle pattern, and lots of brittle fracture areas werepresent on the fracture surface.

(6) After high temperature exposure, the tensile strength of thecopper alloy reduced obviously. The decrease in the tensilestrength makes the bond strength difficult to furtherincrease.

Acknowledgements

The authors gratefully acknowledge the financial support pro-vided by the National Natural Science Foundation of China (No.51205428) and the Fundamental Research Funds for the CentralUniversities (Project No.: CDJZR12130047).

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