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Significantly retarded interfacial reaction between an electroless Ni ...

Significantly retarded interfacial reaction between an electroless Ni–W–P

metallization and lead-free Sn–3.5Ag solder

Ying Yang a , J.N. Balaraju b , Ser Choong Chong c , Hui Xu d , Changqing Liu d , Vadim V. Silberschmidt d ,

Zhong Chen a,⇑

a School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore

b Surface Engineering Division, CSIR National Aerospace Laboratories, Bangalore 560 017, India

c Institute of Microelectronics, A STAR (Agency for Science, Technology and Research), 11 Science Park Road, Singapore Science Park II, Singapore 117685, Singapore

d Wolfson School of Mechanical and Manufacturing Engineering, Loughborough University, Loughborough LE11 3TU, UK

article info

Article history:

Received 22 January 2013

Received in revised form 21 February 2013

Accepted 21 February 2013

Available online 28 February 2013

Keywords:

Metallization

Interfacial reaction

Soldering

Kinetics

Intermetallic compound

Electroless Ni–W–P

1. Introduction

abstract

Due to health and environmental concerns, implementation of

lead-free soldering has commenced since July 1st, 2006. Over the

past two decades, the interfacial reactions between lead-free solders

and Cu-based UBM during reflow and aging as well as their

mechanical properties have been widely studied [1–5]. However,

lead-free solders have higher melting temperatures and higher

Sn content than the conventional eutectic Sn–Pb solder, which

make their reactions with soldering metallization more rapidly,

leading to reliability problems. Thus, it is important to seek solutions

to effectively slow down the interfacial reaction with leadfree

solders.

One way to suppress the interfacial reaction is to change the

solder composition by adding of small quantities of additives into

the solder alloy [6–8], but this method is only applicable to solder

paste. Another approach is to modify the composition of soldering

metallization. Cu-based metallization was found to be unreliable

with lead-free solders as it will be completed reacted away during

reflow process, leaving only intermetallic compound (IMC) at the

⇑ Corresponding author. Tel.: +65 6790 4256.

E-mail address: aszchen@ntu.edu.sg (Z. Chen).

0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved.

http://dx.doi.org/10.1016/j.jallcom.2013.02.113

Journal of Alloys and Compounds 565 (2013) 11–16

Contents lists available at SciVerse ScienceDirect

Journal of Alloys and Compounds

journal homepage: www.elsevier.com/locate/jalcom

To address the potential reliability challenges brought by the accelerated reaction with the adoption of

lead-free solders, an electrolessly-plated Ni–W–P alloy (6–7 wt.% of P and 14–16 wt.% of W) was developed

as the soldering metallization in this study. It was found that the electroless Ni–W–P layer was consumed

much more slowly than the binary Ni–P layer after prolonged reaction. Unlike the Ni–P/Sn–3.5Ag

interface where three intermetallic compounds (IMCs) are formed, only two IMCs are found at the Ni–W–

P/Sn–3.5Ag interface. Besides, there is no void formation at the soldering reaction interface with the ternary

metallization. The growth of Ni3Sn4 and (Ni,W)3P layers at the Ni–W–P/Sn–3.5Ag interface is found

to be diffusion-controlled. The activation energies for the growth of Ni3Sn4 and (Ni,W)3P layers are

62.3 kJ/mol and 58.2 kJ/mol, respectively.

Ó 2013 Elsevier B.V. All rights reserved.

interface. This will cause severe degradation of the mechanical

strength of the joint. It is well known that Ni and its alloys have

a slower reaction rate with Sn than Cu and Cu alloys. Ni-based metallizations

therefore receive great attention with lead-free soldering.

Among Ni-based metallizations, electroless Ni–P has been

extensively studied [9–16]. Soldering reaction between lead-free

solders and electroless Ni–P, in terms of IMC morphology and

growth kinetics has been well understood, so does the joint

strength degradation with thermal treatment [17–19]. However,

formation of Ni 3P and a ternary Ni–Sn–P compound layer at the

reaction interface makes the solder joints more prone to brittle

cracking [17,20]. The industry has adopted electroless Ni–P as soldering

metallization in a number of applications. Nevertheless

from material’s point of view, there is still room to further improve

the reliability of such solder joints by exploring the use of new

materials, which is very important for packages undergoing multiple

reflows.

Incorporation of a third element into electroless Ni–P was

considered as an approach to improve the properties of Ni–P as

a soldering metallization. Ni–W–P alloy is considered as a good

candidate since W is a refractory metal element. Duh’s group

has reported that the crystallization temperature of Ni–P

compounds in electroless Ni–P (8.5 wt.% of P) and electroless

Ni–W–P (7.6 wt.% of P and 10.9 wt.% of W) coatings were 337


12 Y. Yang et al. / Journal of Alloys and Compounds 565 (2013) 11–16

and 406 °C, respectively [21], indicating the addition of W into

Ni–P can effectively retard the crystallization of Ni–P compounds.

Thus, it is possible to slow down the interfacial reaction with

adoption of Ni–W–P as the soldering metallization by retarding

the formation of fast diffusion path, which is grain boundary of

Ni–P compounds in this case. Electroless Ni–W–P film was first

made by Pearlstein and Weightman [22] by adding salt of W into

the electroless nickel bath. Attempts to use the Ni–W–P coating

as the soldering metallization in microelectronic packaging were

first made by Chen et al. [23], and they reported that the

Ni–W–P coating (7 wt.% of P and 14 wt.% of W) with a high W

content have much longer service lifetime than the Ni–P coating

by studying long term interfacial reaction between Ni–W–P and

molten Sn–58Bi solder at 200 °C. After that, the Ni–W–P films

with three different levels of W at two fixed levels of P (5 and

9 wt.%) were prepared by Jang and Yu [24], and the interfacial

reaction between Sn–3.5Ag solder and these Ni–W–P films, as

well as the drop impact test results of these solder joints were

studied. However, the diffusion mechanism and the growth kinetics

of the IMCs at the Ni–W–P/solder interface have not been

reported.

In this work, a Ni–W–P alloy was prepared by electroless plating.

Its interfacial reaction with Sn–3.5Ag solder after reflow and

prolonged aging were studied. The interfacial reaction with a binary

electroless Ni–P under the same reflow and aging conditions

was investigated for comparison. In addition, the growth kinetics

of the Ni 3Sn 4 and (Ni,W) 3P layers was also investigated with their

activation energies reported.

2. Experimental

Cu plates (6 mm thick, 99.98 wt.%) were used as the substrate for both Ni–P and

Ni–W–P plating. Prior to the plating, the Cu substrates were polished and etched

with 35 vol.% nitric acid for 30 s, followed by a commercial Ru activation treatment

for surface activation. Electroless Ni–P plating was carried out in a commercial

acidic sodium hypophosphite bath (from MacDermid) with a pH level of 5.3 (adjusted

by ammonium hydroxide) at 88 ± 2 °C for 40 min. Electroless Ni–W–P plating

was conducted in a self-prepared alkaline bath with a pH level of 9.0

(adjusted by 25% sulphuric acid) at 85 ± 2 °C for 2 h. As listed in Table 1, this alkaline

bath contains nickel sulphate and sodium tungstate as nickel and tungsten

sources, respectively, sodium hypophosphite as a reducing agent along with sodium

citrate as the complexing and buffering agent.

After the plating, Ni–P/Sn–3.5Ag and Ni–W–P/Sn–3.5Ag solder joints were

formed for the study of interfacial reactions. Prior to the soldering, a thin layer of

no-clean paste flux was applied on top of the plated surface of Cu to remove oxides.

Commercially obtained Sn–3.5Ag solder wires with flux in core were used for the

soldering. Solder joining was carried out in an IR reflow oven (ESSEMTEC RO-06E)

with a reflow temperature at 260 °C for 60 s. After the reflow, solid-state aging

was carried out at 200 °C for up to 50 h for the Ni–P/Sn–3.5Ag solder joint and up

to 625 h for the Ni–W–P/Sn–3.5Ag solder joint, respectively. In order to study the

growth kinetics of the IMCs formed at the Ni–W–P/Sn–3.5Ag solder interface, solid-state

aging was also carried out at three lower temperatures (140 °C, 160 °C,

and 180 °C) for up to 625 h.

The surface morphologies and the thicknesses of both deposited layers were observed

by SEM (JEOL JSM-6360A). The interfacial microstructures of both types of

solder joints after prolonged aging at all temperatures were examined using SEM

as well. For cross-sectional SEM, samples were cold mounted in epoxy and polished

down to 1 lm finish, followed by etching with 4% hydrochloric acid to reveal the

interfacial microstructure. The compositions of the as-deposited, the as-reflowed

and the aged samples were analyzed using energy dispersive X-ray (EDX) incorporated

in the SEM.

Table 1

Composition of the plating bath for electroless Ni–W–P.

Constituents of plating bath Concentration (g/L)

NiSO4 6H2O 26

Na2WO4 2H2O 33

NaH2PO2 H2O 12

Na3C6H5O7 2H2O 75

3. Results and discussion

3.1. As-deposited metallizations

The Ni–P coating contains 6–7 wt.% of P. As shown in Fig. 1a, the

surface of the deposited Ni–P layer has smooth nodules with an

uneven size distribution. The thickness was measured to be around

14 lm (Fig. 1b). The Ni–W–P coating contains 6–7 wt.% of P and

14–16 wt.% of W. As shown in Fig. 2a, the surface of the deposited

Ni–W–P layer also contains smooth nodules, but the nodules have

a more even size distribution than the ones in Ni–P. Its thickness

was measured to be around 9.8 lm(Fig. 2b). It was noted that both

coatings have a good adhesion to the Cu substrate (Figs. 1b and 2b).

3.2. Interfacial microstructures after the reflow

The cross-sectional micrograph of as-reflowed Ni–P/Sn–3.5Ag

solder joint is shown in Fig. 3a. A layer of Ni3Sn4 was formed due

to reaction between Sn from the solder and Ni from the metallization.

Some of the faceted Ni3Sn4 particles spalled into the bulk solder.

Beneath the Ni 3Sn 4 layer, there was a very thin layer of Ni 2SnP

[25,26]. Adjacent to the coated Ni–P layer, a dark layer of Ni3P was

present. A number of voids were clearly visible inside the Ni 3P

layer. The formation of these voids has been explained in previous

research [9].

The cross-sectional micrograph of as-reflowed Ni–W–P/Sn–3.5Ag

solder joint is shown in Fig. 3b. It was observed that a layer of

Ni3Sn4 clearly formed, and there was no spallation of IMC as it

did in the binary Ni–P reaction. Under the Ni 3Sn 4 layer, there

was a thin layer consisting of Ni, W, and P elements. The composition

of this ternary layer was measured to be 62.5 at.% of Ni,

10.5 at.% of W and 27 at.% of P by EDX. Such composition is suggestive

of (Ni,W)3P compound. Jang and Yu [24] have conducted the

XRD analysis for the heat-treated Ni–W–P films with different W

levels. The XRD results showed that the Ni3P lattice parameter

did not seem to be affected by the W addition, so they concluded

that W solubility in Ni3P is very limited or the (Ni,W)3P stoichiometry

is quite stable. In addition, this (Ni,W) 3P layer was found to be

amorphous by studying its diffraction pattern under TEM [24]. As

shown in Fig. 3b, the interface between the (Ni,W) 3P layer and

the unconsumed Ni–W–P layer was uneven, implying consumption

of Ni–W–P layer is non-uniform during reflow soldering.

3.3. Solid-state interfacial reactions

Fig. 4a shows the growth of various compounds at the Ni–P/Sn–

3.5Ag interface after aging at 200 °C for 50 h. It was found that

upon 50 h of aging, Ni3Sn4 grew much thicker, and some Ag3Sn

particles accumulated inside this layer. The as-deposited Ni–P

layer was fully consumed and transformed into Ni3P layer. The

thickness of this Ni3P layer ( 7.4 lm) is much smaller than that

of the as-deposited Ni–P layer ( 14 lm). Such shrinkage indicates

that, Ni atoms diffuse out from Ni–P to form Ni 3Sn 4 during the

interfacial reaction. It was observed that only a few voids appeared

in the Ni 3P layer after reflow (Fig. 3a). However, as the reaction

continued upon aging, these voids increased in both the size and

number (Fig. 4a).

Fig. 4b shows the growth of various compounds at the Ni–W–P/

Sn–3.5Ag interfaces after aging at 200 °C for 625 h. It was found

that during aging, the (Ni,W)3P layer grew very slowly. The thickness

increase of the Ni 3Sn 4 layer is much slower than that formed

at the Ni–P/Sn–3.5Ag interface (Fig. 4a). Unlike the Ni–P layer

which was fully consumed after 50 h of aging (Fig. 4a), the Ni

W–P layer after 625 h of aging still had 6.9 lm left, which is

70% of its original thickness. Moreover, it is interesting to note that


even after 625 h aging, no voids were found in the Ni–W–P/Sn–3.5Ag

solder joint. After comparing the two reaction couples, it is

reasonable to attribute the slow IMC growth and consumption rate

of the Ni–W–P metallization pad to the amorphous nature of the

(Ni,W)3P layer which has effectively hindered the diffusion of Ni

from the underneath layer. This point will be elaborated later in

conjunction with diffusion profile analysis.

3.4. IMC growth kinetics at the Ni–W–P/Sn–3.5Ag interface

In the current work, the growth kinetics of both Ni3Sn4 and

(Ni,W) 3P layers was presented. Three aging temperatures (140,

160, and 180 °C) were selected for the study. The average thicknesses

of the Ni 3Sn 4 and (Ni,W) 3P layers were obtained from at

least 10 SEM images of different locations of each sample by measuring

the cross-section area of the layer over a certain length on

the SEM image with the help of an image analyzer. The thickness

of the reaction layer in the solder joint can be generally expressed

by:

Y. Yang et al. / Journal of Alloys and Compounds 565 (2013) 11–16 13

Fig. 1. As-deposited Ni–P layer: (a) surface morphology, (b) cross-sectional micrograph.

Fig. 2. As-deposited Ni–W–P layer: (a) surface morphology, (b) cross-sectional micrograph.

Fig. 3. Back-scattered SEM images showing IMCs formed in the as-reflowed solder joints: (a) Ni–P/Sn–3.5Ag solder joint, and (b) Ni–W–P/Sn–3.5Ag solder joint.

d d0 ¼ kt 1=n

where d and d0 are the thickness of the reaction layer at time t and

zero, respectively, k is the growth rate constant, and n is the time

exponent. Fig. 5 shows the thickness of both the Ni3Sn4 and

(Ni,W) 3P layers as a function of the square root of the aging time

(i.e. assuming n = 2) at different aging temperatures. The thickness

increment of both layers was found to increase linearly with the

square root of aging time, suggesting that the growth of these

two layers at the Ni–W–P/Sn–3.5Ag interface are both controlled

by diffusion. The growth rate constant is calculated from a linear

regression analysis of (d d 0) versus t 0.5 , where the slope is equal

to k, and the values of k for both layers at each aging temperature

are listed in Table 2. The growth rate constants of Ni3Sn4 and

(Ni,W)3P layers increase with increasing aging temperature, indicating

that the growths of both layers were faster at higher aging

temperatures.

To calculate the activation energy for the interfacial compound

growth, the Arrhenius equation is used:

ð1Þ


14 Y. Yang et al. / Journal of Alloys and Compounds 565 (2013) 11–16

(a)

Ag 3Sn

Ni 3P

Voids

Ni 3Sn 4

Ni 2SnP

k 2 ¼ A expð Q=RTÞ ð2Þ

where k 2 is the square of the growth rate constant, A is a prefactor, T

is the absolute temperature, R is the gas constant, and Q is the activation

energy. The value for Q is obtained from the slope of the

Arrhenius plot, as shown in Fig. 6. The activation energy for the

Ni3Sn4 growth in the solid-state reaction between the Ni–W–P

and Sn–3.5Ag solder is obtained to be 62.3 kJ/mol, and the prefactor

is 2.75 10 7 cm 2 /s. Table 3 lists the values of the activation

energies for the growth of Ni3Sn4 in the Ni-based UBM/Sn–3.5Ag

solder systems obtained from previous works [9,27,28]. Our result

lies within the range of values obtained by previous works on

Sn–3.5Ag/Ni or Ni–P reactions. However, when comparing the activation

energies, caution has to be taken that since the solubility of

Ni in liquid Sn is rather high, formation of compound layers actually

takes place under conditions of simultaneous dissolution of the solid

in the melt. The rate of dissolution is known to be dependent on

the experiment geometry, the surface area of contact between the

solid and liquid phases and the volume of the liquid phase in particular.

As a result, compound layer growth kinetics proves to be also

Cu

(b)

(Ni,W)3P

Cu

Ni 3Sn 4

Ni-W-P

Fig. 4. Back-scattered SEM images showing IMCs formed in the aged solder joints: (a) Ni–P/Sn–3.5Ag solder joint after aging at 200 °C for 50 h, and (b) Ni–W–P/Sn–3.5Ag

solder joint after aging at 200 °C for 625 h.

Thickness increment

of Ni3Sn4 (µm)

Thickness increment

of (Ni,W) 3P (µm)

2.4

2

1.6

1.2

0.8

0.4

140 °C

160 °C

180 °C

(a)

0

5 10 15 20 25 30

0.6

0.5

0.4

0.3

0.2

0.1

140 °C

160 °C

180 °C

(b)

Square root of aging time (hour 1/2 )

0

5 10 15 20 25 30

Square root of aging time (hour1/2 )

Fig. 5. Thickness of the (a) Ni 3Sn 4 and (b) (Ni,W) 3P layers formed in the Ni–W–P/

Sn–3.5Ag solder joint during aging at 140, 160, and 180 °C up to 625 h.

Table 2

IMC growth rate constants at the Ni–W–P/Sn–3.5Ag interface at various aging

temperatures.

Temperatures (°C) k of Ni 3Sn 4

( 10 8 cm/s 1/2 )

140 6.13 2.13

160 8.78 2.20

180 13.7 2.30

k of (Ni,W) 3P

( 10 8 cm/s 1/2 )

dependent on the experiment geometry. Since it is usually different

in different works, close values of the activation energy can hardly

be expected. Similarly, the reaction kinetics in solid-state reaction

could be affected by the difference in these experimental conditions

too. The twofold difference observed in its values as listed in Table 3

could be due to this reason.

Despite the possible discrepancy, the activation energy is an

important parameter as it reflects the magnitude of the critical barrier

for the compound growth. The comparable activation energy

value with other reported Sn–3.5Ag/Ni-based soldering systems

indicates that controlling mechanism for the Ni3Sn4 IMC formation

ln (k 2 in cm 2 /s) of Ni 3 Sn 4

ln (k 2 in cm 2 /s) of (Ni,W) 3 P

-31.2

-31.6

-32

-32.4

-32.8

-33.2

(a)

-33.6

2.15 2.2 2.25 2.3

1000/T (1/K)

2.35 2.4 2.45

-35.15

-35.2

-35.25

-35.3

-35.35

(b)

-35.4

2.15 2.2 2.25 2.3 2.35 2.4 2.45

1000/T (1/K)

Fig. 6. Arrhenius plot of the growth of (a) Ni3Sn4 and (b) (Ni,W)3P layers in the Ni

W–P/Sn–3.5Ag solder joint.


Table 3

Activation energy for the growth of Ni 3Sn 4 in the Ni-based UBM/Sn–3.5Ag solder

systems obtained in different studies.

Solder/substrate Experimental conditions

(temperature/time)

Q (kJ/mol) Reference

Sn–3.5Ag/Ni–W–P 140–180 °C/up to 625 h 62.3 Present work

Sn–3.5Ag/Ni–P 130–170 °C/up to 625 h 110.0 [9]

Sn–3.5Ag/Ni–P 140–200 °C/up to 400 h 98.9 [26]

Sn–3.5Ag/Ni–P 100–170 °C/up to 60 days 49.0 [27]

Sn–3.5Ag/Au/Ni 70–170 °C/up to 100 days 72.5 [28]

Fig. 7. Elemental distribution at Ni–W–P/Sn–3.5Ag interface after aging at 200 °C

for 225 h.

might remain the same, which is likely to be the diffusion of Sn in

the Ni 3Sn 4 IMC. As cited in a review article by Ho et al. [29], Ti marker

experiment in a Ni/Sn reaction couple has proven that Sn diffuses

faster than Ni through Ni3Sn4 IMC, therefore the IMC

growth activation energy is likely related to the Sn diffusivities in

Y. Yang et al. / Journal of Alloys and Compounds 565 (2013) 11–16 15

the Ni3Sn4 IMC. Unfortunately, data for Sn diffusivities in Ni3Sn4

are unavailable, which could be an interesting topic for future

studies.

The activation energy for the (Ni,W) 3P growth in the solid-state

solder reaction is found to be 58.2 kJ/mol, and the prefactor is

2.46 10 15 cm 2 /s. It is worth noting that the prefactors for

growth of both the Ni3Sn4 and (Ni,W)3P layers are many orders

of magnitude lower than the prefactor of Ni 3Sn 4 formation in soldering

with Ni–P [9], which is due to the presence of W in the metallization

pad that decreases the availability of Ni at the soldering

interface. Although the activation energy of (Ni,W)3P growth is

similar to that of Ni 3Sn 4 growth, the prefactor of (Ni,W) 3P growth

is much lower, so (Ni,W)3P grew much slower than Ni3Sn4.

3.5. Formation of IMCs at the Ni–W–P/Sn–3.5Ag interface

EDX line scan and element mapping analysis were obtained to

understand the diffusion profile at the Ni–W–P/Sn–3.5Ag interface.

Fig. 7 shows elemental distribution at Ni–W–P/Sn–3.5Ag interface

after aging at 200 °C for 225 h. It was noted that at the interface between

the Ni–W–P layer and the (Ni,W) 3P layer, Ni intensity

started to decrease. This indicates that Ni atoms diffuse outward

from the deposited Ni–W–P layer during the interfacial reaction,

leading to the formation of Ni3Sn4. A sudden drop of Sn intensity

was observed at the (Ni,W) 3P/Ni 3Sn 4 interface, which indicates

there is no Sn in the (Ni,W)3P layer.

From the mapping result of the sample after aging at 200 °C for

625 h (Fig. 8), it is clearly seen that W and P remains in the Ni–W–P

and (Ni,W) 3P layers. In contrast, Ni atoms diffuse out from the

Ni–W–P layer to form Ni3Sn4. The signal from Sn was found to

Fig. 8. EDX element mapping analysis of the Ni–W–P/Sn–3.5Ag interface after aging at 200 °C for 625 h: (a) SEM image, (b) mapping for Ni, (c) mapping for W, (d) mapping

for P, (e) mapping for Sn, (f) mapping for Cu, and (g) mapping for Ag. Elemental concentration decreases with increasing black intensity.


16 Y. Yang et al. / Journal of Alloys and Compounds 565 (2013) 11–16

terminate at the Ni 3Sn 4/(Ni,W) 3P interface, and no signal from Sn

was detected in the (Ni,W)3P layer, indicating that Sn atoms did

not diffuse through the (Ni,W) 3P layer. These are in agreement

with the findings from the line scan (Fig. 7). It was also noted that

after such a prolonged aging, Cu atoms still remain in the substrate,

and no (Cu,Ni)6Sn5 or (Ni,Cu)3Sn4 compounds formed, suggesting

that Ni–W–P is a good diffusion barrier preventing inter-diffusion

between Cu and Sn.

Fig. 9 illustrates the diffusional formation mechanism of the

Ni3Sn4 and (Ni,W)3P layers at the Ni–W–P/Sn–3.5Ag interface

schematically. Based on elemental analysis, P atoms in the Ni

W–P remained in the deposited layer during soldering reaction.

The deposited Ni–W–P layer contains 82 at.% of Ni, 5 at.% of W

and 13 at.% of P, while the newly-formed (Ni,W)3P layer contains

62.5 at.% of Ni, 10.5 at.% of W and 27 at.% of P. Thus, the ratio of

W to P in the (Ni,W)3P layer ( 2.6) is the same as that in the

deposited Ni–W–P layer, suggesting that Ni was the only element

to diffuse outward from the Ni–W–P coating layer during soldering

reaction. The driving force for the Ni diffusion is the Ni concentration

difference between the Ni–W–P coating layer and the bulk solder,

which contains no Ni. As shown in Fig. 9, since concentration

of Ni is higher in the coating layer, Ni atoms diffuse out from the

Ni–W–P layer towards solder and react with Sn atoms from the

solder to form Ni3Sn4. Meanwhile, the formation of Ni3Sn4 causes

the depletion of Ni from the surface of the Ni–W–P layer, leading

to continued growth of the newly formed ternary (Ni,W)3P layer.

It was reported that the Ni3P layer formed at the Ni–P/Sn–3.5Ag

interface has a fine columnar structure, so this layer acts as a fast

diffusion path for Ni atoms during the interfacial reaction [9]. As

a result, the rapid diffusion of Ni atoms through the Ni 3P layer

leads to the formation of voids inside this layer while there is

not enough compensation from other elements to fill the vacant

sites left by Ni. Jang and Yu have reported that the (Ni,W)3P layer

formed at the Ni–W–P/Sn–3.5Ag interface has an amorphous

structure [24]. Since the amorphous structure of this (Ni,W)3P

layer is free of fast diffusion channel such as grain boundaries, it

is more difficult for Ni atoms to diffuse out through this layer. This

is evidenced by observation that no voids are formed inside this

layer after prolonged reaction (Fig. 4b). As a result, the interfacial

IMC growth rate at the Ni–W–P/Sn–3.5Ag interface was much

slower than that at the Ni–P/Sn–3.5Ag interface. This is also evidenced

by the much slower consumption rate of the Ni–W–P metallization

compared to the binary Ni–P.

4. Conclusions

Sn-3.5Ag solder

Ni 3 Sn 4

(Ni,W) 3 P

Electroless Ni-W-P

Cu Surface

Fig. 9. A simplified scheme illustrating the diffusional formation mechanism of

Ni3Sn4 and (Ni,W)3P layers between Ni–W–P metallization and Sn–3.5Ag solder.

In this work, an electrolessly-plated Ni–W–P alloy (6–7 wt.% of

P and 14–16 wt.% of W) was developed as an alternative Ni-based

Sn

Ni

metallization for lead-free soldering. Interfacial reaction between

Sn–3.5Ag solder and electroless Ni–W–P after reflow and prolonged

aging was investigated, with interfacial reaction between

the same solder and electroless Ni–P (6–7 wt.% of P) as a benchmark.

At the Ni–W–P/Sn–3.5Ag interface, only two interfacial compounds,

Ni 3Sn 4 and (Ni,W) 3P are formed, and the diffusion

mechanism was proposed. After prolonged aging, unlike at the

Ni–P/Sn–3.5Ag interface, no voids were found at the Ni–W–P/Sn–

3.5Ag interface, and the Ni3Sn4 grew much more slowly. These

facts indicate a significantly slow out-diffusion of Ni atoms at the

Ni–W–P/Sn–3.5Ag interface; hence, Ni–W–P layer stands out as

an impressive diffusion barrier for lead-free soldering. In addition,

the growth kinetics of the Ni3Sn4 and (Ni,W)3P layers at the Ni–W–

P/Sn–3.5Ag interface was examined as a consequence of aging at

three temperatures (140, 160 and 180 °C) for up to 625 h. The

growth of these two layers was found to be a diffusion-controlled

process. The activation energies for the growth of Ni3Sn4 and

(Ni,W)3P layers during solid-state reaction are 62.3 kJ/mol and

58.2 kJ/mol, respectively.

Acknowledgements

The authors very much appreciate technical discussions with

Dr. K. Chen and Prof. K.N. Tu. Financial assistance from MOE Singapore

(Grant RG 19/00, RG 14/03), and UK Department for Innovation,

Universities and Skills (DIUS) through a PMI2 Project (Grant

No. RC 41) is gratefully acknowledged.

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