U.S. patent application number 13/066748 was filed with the patent office on 2011-12-15 for pb-free sn-ag-cu-al or sn-cu-al solder.
This patent application is currently assigned to Iowa State University Research Foundation, Inc.. Invention is credited to Iver E. Anderson, Adam J. Boesenberg, Joel L. Harringa.
Application Number | 20110303448 13/066748 |
Document ID | / |
Family ID | 45095312 |
Filed Date | 2011-12-15 |
United States Patent
Application |
20110303448 |
Kind Code |
A1 |
Anderson; Iver E. ; et
al. |
December 15, 2011 |
Pb-Free Sn-Ag-Cu-Al or Sn-Cu-Al Solder
Abstract
A solder alloy includes Sn, optional Ag, Cu, and Al wherein the
alloy composition is controlled to provide a strong, impact-and
thermal aging-resistant solder joint that has beneficial
microstructural features and is substantially devoid of Ag.sub.3Sn
blades.
Inventors: |
Anderson; Iver E.; (Ames,
IA) ; Harringa; Joel L.; (Ames, IA) ;
Boesenberg; Adam J.; (Ames, IA) |
Assignee: |
Iowa State University Research
Foundation, Inc.
|
Family ID: |
45095312 |
Appl. No.: |
13/066748 |
Filed: |
April 22, 2011 |
Related U.S. Patent Documents
|
|
|
|
|
|
Application
Number |
Filing Date |
Patent Number |
|
|
61343135 |
Apr 23, 2010 |
|
|
|
Current U.S.
Class: |
174/259 ;
174/84R; 228/101; 228/56.3; 420/560 |
Current CPC
Class: |
B23K 35/0244 20130101;
B23K 35/262 20130101; H05K 3/3463 20130101; B23K 35/001 20130101;
C22C 13/00 20130101 |
Class at
Publication: |
174/259 ;
420/560; 228/101; 228/56.3; 174/84.R |
International
Class: |
H05K 1/02 20060101
H05K001/02; H01R 4/02 20060101 H01R004/02; B23K 35/14 20060101
B23K035/14; C22C 13/00 20060101 C22C013/00; B23K 31/02 20060101
B23K031/02 |
Goverment Interests
CONTRACTUAL ORIGIN OF THE INVENTION
[0002] This invention was made with government support under
Contract No. DE-ACO2-07CH11358 awarded by the U.S. Department of
Energy. The government has certain rights in the invention.
Claims
1. An alloy consisting essentially of about 3 to about 4 weight %
Ag, about 0.7 to about 1.7 weight % Cu, about 0.01 to about 0.25
weight % Al, and balance consisting essentially of Sn.
2. The alloy of claim 1 having a solidus temperature of about
217.degree. C.
3. The alloy of claim 2 having a liquid plus solid temperature
range less than about 5 degrees C.
4. The alloy of claim 1 having tin dendrites, a ternary eutectic
between the dendrites, and pro-eutectic particles adjacent and/or
within the tin dendrites in an as-solidified microstructure.
5. The alloy of claim 4 wherein the pro-eutectic particles comprise
Cu.sub.6Sn.sub.5.
6. The alloy of claim 4 wherein the alloy is as-solidified at less
than 5.degree. C./second.
7. The alloy of claim 1 having an Al content selected to form
particles in the solder that comprise Cu and Al.
8. The alloy of claim 7 wherein the particles comprises
substantially Cu.sub.33Al.sub.17.
9. A solder alloy consisting essentially of about 3.4 to about 3.6
weight % Ag, about 0.8 to about 1.1 weight % Cu, about 0.03 to
about 0.20 weight % Al, and balance consisting essentially of
Sn.
10. A solder alloy consisting essentially of about 3.45 to about
3.55 weight % Ag, about 0.75 to about 1.0 weight % Cu, about 0.04
to about 0.15 weight % Al, and balance consisting essentially of
Sn.
11. A solder alloy consisting essentially of about 3 to about 4
weight % Ag, 0.95-y weight % Cu, and y weight % Al and balance
consisting essentially of Sn wherein y is about 0.01 to about 0.25
weight %.
12. A solder alloy comprising Sn, Ag, Cu, and Al and having a
solidus temperature of about 217.degree. C. and a narrow
liquid-solid mushy zone with a liquidus temperature not exceeding
about 5.degree. C. above the solidus temperature.
13. A solder alloy for BGA applications, consisting essentially of
about 3.45 to about 3.55 weight % Ag, about 0.80 to about 1.0
weight % Cu, about 0.10 to about 0.20 weight % Al, and balance
consisting essentially of Sn.
14. A solder alloy consisting essentially of about 0.7 to about 3.5
weight % Cu, about 0.01 to about 0.25 weight % Al, and balance
consisting essentially of Sn.
15. The solder alloy of claim 11 consisting essentially of about
0.8 to about 3.2 weight % Cu, about 0.03 to about 0.25 weight % Al,
and balance consisting essentially of Sn.
16. The alloy of claim 11 consisting essentially of about 0.95 to
about 3.0 weight % Cu, about 0.15 to about 0.20 weight % Al, and
balance consisting essentially of Sn.
17. A solder alloy consisting essentially of about 3.20-y weight %
Cu, and y weight % Al and balance consisting essentially of Sn
wherein y is about 0.15 to about 0.25 weight %.
18. A solder joint comprisng a Sn--Ag--Cu--Al solder alloy
soldified in contact with an electrical conductor wherein the
solder joint has an as-solidified microstructure that comprises tin
dendrties, ternary eutectic between tin dendrites, and pro-eutectic
Cu.sub.6Sn.sub.5 particles adjacent and/or within the tin dendrites
and that is substantially devoid of Ag.sub.3Sn blades.
19. The joint of claim 18 having an interfacial layer comprising
Cu.sub.6Sn.sub.5 and an adjacent metastable intermediate rejected
solute region as a zone of intermediate hardness between the
interfacial layer and the solder tin matrix.
20. The joint of claim 18 wherein the solder alloy consists
essentially of about 3 to about 4 weight % Ag, about 0.7 to about
1.7 weight % Cu, about 0.01 to about 0.25 weight % Al, and balance
consisting essentially of Sn.
21. The joint of claim 20 wherein the solder alloy consists
essentially of about 3.4 to about 3.6 weight % Ag, about 0.8 to
about 1.1 weight % Cu, about 0.03 to about 0.20 weight % Al, and
balance consisting essentially of Sn.
22. The joint of claim 21 wherein the solder alloy consists
essentially of about 3.45 to about 3.55 weight % Ag, about 0.75 to
about 1.0 weight % Cu, about 0.04 to about 0.15 weight % Al, and
balance consisting essentially of Sn.
23. The joint of claim 18 formed on an electrical wiring board.
24. The joint of claim 18 formed about copper electrical
conductors.
25. The joint of claim 18 having particles in and/or adjacent a
solder joint interface layer wherein the particles comprise Cu and
Al.
26. The joint of claim 25 wherein the particles comprises
substantially Cu.sub.33Al.sub.17.
27. A thermally-aged solder joint comprising a Sn--Ag--Cu--Al
solder alloy in contact with an electrical conductor wherein the
solder joint has an interfacial layer thickness that is about the
same as the thickness of the interfacial layer in the as-solidified
solder before thermal aging.
28. The joint of claim 27 wherein the interfacial layer thickness
is no more than 30% greater than the thickness of the interfacial
layer in the as-solidified solder before thermal aging.
29. The joint of claim 27 wherein the solder alloy consists
essentially of about 3 to about 4 weight % Ag, about 0.7 to about
1.7 weight % Cu, about 0.01 to about 0.25 weight % Al, and balance
consisting essentially of Sn.
30. A solder joint comprisng a Sn--Ag--Cu--Al solder alloy
soldified in contact with an electrical conductor wherein the
solder joint has an as-solidified microstructure that comprises an
interfacial layer comprising Cu.sub.6Sn.sub.5 and an adjacent
metastable intermediate Al-containing rejected solute region
between the interfacial layer and the solder tin matrix, wherein
the hardness of the Al-containing rejected solute region is
intermediate the hardness of the interfacial layer and the solder
tin matrix.
31. The joint of claim 30 formed on an electrical wiring board.
32. The joint of claim 30 formed about copper electrical
conductors.
33. The joint of claim 30 wherein the solder alloy consists
essentially of about 3 to about 4 weight % Ag, about 0.7 to about
1.7 weight % Cu, about 0.01 to about 0.25 weight % Al, and balance
consisting essentially of Sn.
34. The joint of claim 33 wherein the solder alloy consists
essentially of about 3.4 to about 3.6 weight % Ag, about 0.8 to
about 1.1 weight % Cu, and about 0.03 to about 0.20 weight % Al,
and balance consisting essentially of Sn.
35. The joint of claim 34 wherein the solder alloy consists
essentially of about 3.45 to about 3.55 weight % Ag, about 0.75 to
about 1.0 weight % Cu, and about 0.04 to about 0.15 weight % Al,
and balance consisting essentially of Sn.
36. The joint of claim 30 having particles in and/or adjacent the
solder joint interfacial layer wherein the particles comprise Cu
and Al.
37. The joint of claim 36 wherein the particles comprises
substantially Cu.sub.33Al.sub.17.
38. In a soldering process, the step of solidifying a molten
Pb-free solder consisting essentially of about 3 to about 4 weight
% Ag, about 0.7 to about 1.7 weight % Cu, and about 0.01 to about
0.25 weight % Al, and balance consisting essentially of Sn.
39. The process of claim 38 wherein the solder is solidified on an
electrical wiring board.
40. The process of claim 38 wherein the solder is solidified about
copper electrical conductors.
41. The process of claim 38 including forming buoyant particles
comprising Cu and Al in the molten solder.
42. In a solder paste reflow or BGA solder process, the step of
solidifying a molten Pb-free solder consisting essentially of about
3 to about 4 weight % Ag, about 0.7 to about 1.7 weight % Cu, and
about 0.01 to about 0.25 weight % Al, and balance consisting
essentially of Sn.
43. The process of claim 42 wherein the solder is cooled at a rate
to form an as-solidified microstructure that comprises tin
dendrites, ternary eutectic between the tin dendrites, and
pro-eutectic Cu.sub.6Sn.sub.5 particles adjacent and/or within the
tin dendrites and that is substantially devoid of Ag.sub.3Sn
blades.
44. The process of claim 42 that forms an interfacial layer
comprising Cu.sub.6Sn.sub.5 and an adjacent metalstable
intermediate rejected solute region as a zone of intermediate
hardness between the interfacial layer and the solder tin
matrix.
45. The process of claim 42 wherein the solder is solidified on an
electrical wiring board.
46. The process of claim 42 wherein the solder is solidified about
copper electrical conductors.
47. The process of claim 42 wherein the solder alloy consists
essentially of about 3.45 to about 3.55 weight % Ag, about 0.80 to
about 1.0 weight % Cu, about 0.10 to about 0.20 weight % Al, and
balance consisting essentially of Sn.
48. A solder ball comprising an alloy consisting essentially of
about 3 to about 4 weight % Ag, about 0.7 to about 1.7 weight % Cu,
about 0.01 to about 0.25 weight % Al, and balance consisting
essentially of Sn.
49. The solder ball of claim 48 having a liquid plus solid
temperature range less than. about 5 degrees C.
50. The solder ball of claim 48 wherein the alloy consists
essentially of about 3.4 to about 3.6 weight % Ag, about 0.8 to
about 1.1 weight % Cu, about 0.03 to about 0.20 weight % Al, and
balance consisting essentially of Sn.
51. The solder ball of claim 50 wherein the solder alloy consists
essentially of about 3.45 to about 3.55 weight % Ag, about 0.75 to
about 1.0 weight % Cu, about 0.04 to about 0.15 weight % Al, and
balance consisting essentially of Sn.
52. The solder ball of claim 50 wherein the solder alloy consists
essentially of about 3.45 to about 3.55 weight % Ag, about 0.80 to
about 1.0 weight % Cu, about 0.10 to about 0.20 weight % Al, and
balance consisting essentially of Sn.
Description
[0001] This application claims benefits and priority of U.S.
provisional application Ser. No. 61/343,135 filed Apr. 23, 2010,
the disclosure of which is incorporated herein by reference.
FIELD OF THE INVENTION
[0003] The present invention provides a Pb-free solder alloy
(Sn--Ag--Cu--Al or Sn--Cu--Al) that displays reliable joint
solidification control to provide a strong, impact- and thermal
aging-resistant solder joint having beneficial microstructural
features and substantially devoid of Ag.sub.3Sn blades and that is
useful for joining electronic assemblies and electrical contacts
and to substitute for Pb-containing solders in all surface mount
solder assembly operations, including solder paste reflow and ball
grid array joints.
BACKGROUND OF THE INVENTION
[0004] The global electronic assembly community is striving to
accommodate the replacement of Pb-containing solders, primarily
Sn--Pb alloys, with Pb-free solders due to environmental
regulations and market pressures. During this major transition away
from eutectic or near-eutectic Sn--Pb solder (T.sub.eut=183.degree.
C.) for electronic assembly, there is also the opportunity to make
a major improvement in Pb-free joint reliability for challenging
operating environments, i.e., high temperatures and stress levels,
as well as impact loading situations. Of the Pb-free choices, an
array of solder alloys based on the Sn3.7Ag-0.9Cu (in wt. %)
ternary eutectic (T.sub.eut=217.degree. C.) composition have
emerged with the most potential for broad use across the industry.
U.S. Pat. No. 5,527,628 describes such Pb-free solder alloys.
[0005] The electronics industry has seized the challenge of
adaptation and is proceeding rapidly to develop the assembly
techniques and to generate the reliability data for
tin-silver-copper (SAC) alloy solder as a favored Pb-free solder in
many electronic assembly applications. Compared with Sn--Pb solders
that have been limited typically to low-stress joints and
reduced-temperature service because of the soft Pb phase that is
prone to coarsening and ductile creep failure, the high Sn content
and strong intermetallic phases of a well-designed SAC alloy solder
can promote enhanced joint shear strength and creep resistance and
can permit an increased operating temperature envelope for advanced
electronic systems and devices.
[0006] Results of SAC alloy development have demonstrated increased
shear strength at ambient temperature and elevated temperatures,
e.g., 150.degree. C. Joints made from a variety of SAC solders have
also demonstrated resistance to isothermal fatigue and resistance
to degradation of shear strength from thermal aging for temperature
excursions up to 150.degree. C., a current test standard for
under-the-hood automotive electronics solder.
[0007] An observation that arose from initial widespread testing of
SAC solder alloys was the occasional embrittlement of SAC solder
joints due to micro-void nucleation, growth, and coalescence, if
the exposure to elevated temperatures was sufficiently high,
typically greater than about 150.degree. C., and the exposure was
sufficiently long, greater than about 500 to 1000 h (hours). This
occasional joint embrittlement after thermal aging was observed at
elevated Cu content in SAC solder alloys and typically was
associated with excessive growth of layers of Cu-base intermetallic
compounds, Cu.sub.6Sn.sub.5 and, especially, Cu.sub.3Sn. It should
be noted that U.S. Pat. No. 6,231,691 provides a solder to suppress
this thermal aging phenomenon through minor additions (<1 wt. %,
but usually 0.2-0.3 wt. %) of a fourth element, such as Ni, Fe,
and/or Co, and "like-acting elements," to the SAC solder to
suppress solid state diffusion at the solder/substrate interface
that contains the Cu-base intermetallic compound layers. Later
testing showed that a Mn addition was one of the most effective
like-acting elemental additions, suppressing growth of both types
of intermetallic layers after extensive thermal aging. This type of
minor alloy addition to prevent embrittlement has become
increasingly important since narrow solder joint gaps are becoming
more common with miniaturization of electronic circuits.
[0008] Studies have shown that Sn dendrites are the dominant
as-solidified microstructure feature in solder joints made with
many SAC alloys, not a fine (ternary) eutectic, contrary to the
previous experience with Sn--Pb. Also, it was found that relatively
high but variable undercooling was observed commonly before joint
solidification leading to Sn dendrites with spacing variations
(that depend on undercooling and growth rate) but with very few
distinct Sn grains. The unusually high undercooling of the SAC
solder joints was associated with the difficulty of nucleating Sn
solidification, as a pro-eutectic phase. Especially during slow
cooling, e.g., in ball grid array (BGA) joints where cooling rates
are less than 0.2.degree. C./s, increased undercooling of the
joints also can promote formation of undesirable pro-eutectic
intermetallic phases, specifically Ag.sub.3Sn "blades," that tend
to coarsen radically, leading to embrittlement of as-solidified
solder joints.
[0009] References 1, 2, 3, and 4 listed below employed fourth
element additions to SAC solders with the intention of avoiding
Ag.sub.3Sn blades by selecting SAC compositions that were deficient
in Ag and Cu, e.g., see SAC2705 [see ref. 4], SAC305, and SAC 105
[see refs. 1,2,3]. These references include the following:
[0010] 1. A. W. Liu and N-C. Lee, "The Effects of Additives to
SnAgCu Alloys on Microstructure and Drop Impact Reliability of
Solder Joints," JOM, 59, no. 7 (2007) pp. 26-31.
[0011] 2. B. L-W. Lin et al., "Alloying modification of Sn--Ag--Cu
solders by manganese and titanium," Microelectron. Reliab. (2008),
doi:10.1016/j.microre1.2008.10.001.
[0012] 3. C. W. Liu, P. Bachorik, and N-C. Lee, "The Superior Drop
Test Performance of SAC-Ti Solders and Its' Mechanism," Proc 58th
Electronic Components and Technology Conf, (2008), pp. 627-635.
[0013] 4. D. S. K. Kang, P. A. Lauro, D.-Y. Shih, D. W. Henderson,
K. J. Puttlitz, IBM J. Res. Dev. 49(4/5), 607-620 (2005).
[0014] In these references, some marginally near-eutectic SAC alloy
designs were proposed with a low Cu level (0.5%) and very low Ag
levels, less than 2.7% Ag [ref 4] and down to 1% Ag (SAC 105).
These base alloys were selected since they would tend to promote Sn
formation and inhibit nucleation of Ag.sub.3Sn [ref. 1, 2, 3, 4].
Because these alloys deviate increasingly from the eutectic, they
exhibit a wider melting range (mushy zone) with a liquidus
temperature (for SAC 105) as high as 226.degree. C. Unfortunately,
some observations of unmodified SAC 105 interfacial failure on
impact loading still occurred, since occasional high undercooling
still may permit Ag.sub.3Sn blade formation during slow cooling.
These "interfacial adhesion" failures prompted attempts at alloy
modifications of SAC 105 solder with 1-2 additions [refs. 1,2] to
improve impact resistance of BGA joints by increasing the
interfacial bond strength of the intermetallic layer and,
presumably, by suppressing Ag.sub.3Sn blade formation. While
significant improvement in impact resistance was observed,
especially for SAC105+0.13% Mn and SAC105+0.02% Ti alloys [ref. 3]
(and no Ag3Sn blades were reported), their high liquidus
temperature (approximately 226 .degree. C.). and wide liquid-solid
mushy zone (equal to 9.degree. C. because of the 217.degree. C.
solidus temperature) inhibits broad service application.
SUMMARY OF THE INVENTION
[0015] In an embodiment, the present invention provides a solder
alloy comprising Sn, optional Ag, Cu, and Al wherein the alloy
composition is controlled to provide a strong, impact- and thermal
aging-resistant solder joint having beneficial microstructural
features and substantially devoid of Ag.sub.3Sn blades.
[0016] An illustrative embodiment of the invention provides a
solder alloy consisting essentially of about 3 to about 4 weight %
Ag, about 0.7 to about 1.7 weight % Cu, about 0.01 to about 0.25
weight % Al, and balance consisting essentially of Sn. A preferred
solder alloy consists essentially of about 3.4 to about 3.6 weight
% Ag, about 0.8 to about 1.1 weight % Cu, about 0.03 to about 0.20
weight % Al, and balance consisting essentially of Sn. A still more
preferred solder alloy consists essentially of about 3.45 to about
3.55 weight % Ag, about 0.9 to about 1.0 weight % Cu, about 0.04 to
about 0.10 weight % Al, and balance consisting essentially of Sn.
Another still more preferred solder alloy consists essentially of
about 3.45 to about 3.55 weight % Ag, about 0.75 to about 1.0
weight % Cu, about 0.04 to about 0.15 weight % Al, and balance
consisting essentially of Sn.
[0017] Another illustrative embodiment of the invention provides a
solder alloy consisting essentially of about 3 to about 4 weight %
Ag, 0.95-y weight % Cu, and y weight % Al and balance consisting
essentially of Sn wherein y is about 0.01 to about 0.25 weight
%.
[0018] Still another embodiment of the invention provides a still
more preferred solder alloy consists essentially of about 3.45 to
about 3.55 weight % Ag, about 0.80 to about 1.0 weight % Cu, about
0.10 to about 0.20 weight % Al, and balance consisting essentially
of Sn, especially for BGA applications that involve
thermal-mechanical fatigue environments, like avionics.
[0019] The invention also envisions a modification of the alloy
formulation to eliminate the Ag component for situations where
higher solder melting alloys can be tolerated. Such modified
embodiment provides a solder alloy consisting essentially of about
0.7 to about 3.5 weight % Cu, about 0.01 to about 0.25 weight % Al,
and balance consisting essentially of Sn. A preferred solder alloy
consists essentially of about 0.8 to about 3.2 weight % Cu, about
0.03 to about 0.25 weight % Al, and balance consisting essentially
of Sn. A still more preferred embodiment of this solder alloy
consists essentially of about 0.95 to about 3.0 weight % Cu, about
0.15 to about 0.20 weight % Al, and balance consisting essentially
of Sn.
[0020] Another such modified embodiment involves an the alloy
formulation to eliminate the Ag component for situations where
higher solder melting alloys can be tolerated, where another
illustrative embodiment of the invention provides a solder alloy
consisting essentially of about 3.20-y weight % Cu, and y weight %
Al and balance consisting essentially of Sn wherein y is about 0.15
to about 0.25 weight %.
[0021] The present invention also provides a solder joint and
solder process that embody a Sn-optional Ag--Cu--Al alloy of the
type discussed above. The solder joint has a microstructure that
comprises tin dendrites, interdendritic multi-phase ternary
eutectic, and pro-eutectic Cu.sub.6Sn.sub.5 particles adjacent
and/or within the tin dendrites and is substantially devoid of
Ag.sub.3Sn blades. The as-solidified solder joint microstructure
typically includes an interfacial layer comprising Cu.sub.6Sn.sub.5
and an adjacent metastable Al-containing rejected solute region as
a zone of intermediate hardness between the hard, brittle
interfacial layer and the softer tin matrix of the solder
microstructure to provide a beneficial hardness gradient that
improves impact resistance. The interfacial layer resides between
the copper substrate and the tin matrix. In an embodiment of the
invention, relatively hard particles comprising Cu and Al, such as
Cu.sub.33Al.sub.17, can be formed at the interfacial layer (i.e. in
and/or adjacent the interfacial layer). The solder joint is formed
by the solder being solidified on an electrical wiring board and/or
about copper electrical conductors in illustrative embodiments of
the invention.
[0022] The thermally-aged solder joint has an interfacial layer
thickness that is about the same as the thickness as the
interfacial thickness in the as-solidified condition (e.g. no more
than 30% greater in thickness).
[0023] Solder joints made with the Sn-optional Ag--Cu--Al solder
alloy may need to accommodate some minor addition of Pb due to
reflow and mixing with Sn--Pb component lead plating during reflow
assembly of solder joints. Slight contamination by such small Pb
levels is not expected to raise the beneficial (about 217.degree.
C.) melting point of the Sn--Ag--Cu--AI solder alloys of the
invention and may even help improve the wettability during joint
formation. This type of Pb-tolerant behavior is an advantage over
competing Sn--Ag--Bi (Pb-free) solders that run the risk of
generating an extremely low melting Sn--Pb--Bi ternary eutectic, if
alloyed with Sn--Pb component platings. It is expected that the
global supply of "legacy" electronic components with Sn--Pb solder
plating will continue to diminish and eventually vanish during the
current transition to full Pb-free electronic soldering, but this
possibility must be tolerated in any new Pb-free solders that are
proposed.
[0024] The Sn--Ag--Cu--Al solder alloy exhibits a reduced melting
point (solidus or melting temperature) of about 217.degree. C. as
compared to the melting points of Sn--Ag binary eutectic solder
(221.degree. C.), and the Sn--Cu binary eutectic solder
(227.degree. C.) and a narrow liquid-solid mushy zone with a
liquidus temperature not exceeding about 5.degree. C., preferably
3.degree. C., above the solidus temperature for solderability. This
significantly reduced melting point is a great advantage for solder
assembly of electronic circuits and electrical systems. In the type
of solder paste reflow and ball grid array (BGA) applications that
are envisioned for use with the Sn--Ag--Cu--Al solder, every single
degree of reduced reflow temperature is a precious advantage for
reducing damage to temperature sensitive electronic components and
to the circuit board material, itself. In fact, a reason that SAC
solder came into broad use as a Pb-free alternative to Sn--Pb
solder is that the minimum reflow temperature of SAC solder for
most applications, about 240.degree. C., is just below the
threshold for significant damage of one of the most popular circuit
board materials, a fiberglass/epoxy composite, i.e., FR-4. Thus,
the Sn--Ag--Cu--Al solder alloy pursuant to the present invention
should permit a more comfortable margin for preventing thermal
damage of most components and common circuit board materials.
[0025] The above advantages of the invention will become more
readily apparent from the following detailed description taken with
the following drawings.
BRIEF DESCRIPTION OF THE DRAWINGS
[0026] FIG. 1 is bar graph illustrating the effect of aluminum
additions (in weight %) to SAC 3595 solder alloy on undercooling
values for DSC (differential scanning calorimetry) measurements at
0.17.degree. C./second cooling rate.
[0027] FIG. 2 shows a summary of DSC results for SAC3595+0.01% Al,
SAC3595+0.025% Al, and SAC3595+0.05% Al solder joints (where % is
weight %).
[0028] FIGS. 3a and 3b are photomicrographs of the as-solidified
microstructure of the SAC3595+0.01% Al solder joint and the
SAC3595+0.025% Al solder joint, respectively, cooled at
0.17.degree. C./s in selected DSC tests wherein the as-solidified
microstructure comprises tin dendrites, interdendritic ternary
eutectic, and pro-eutectic Cu.sub.6Sn.sub.5 particles adjacent
and/or within the tin dendrites and the microstructure is devoid of
Ag.sub.3Sn blades.
[0029] FIG. 3c is an SEM of the typical as-solidified
microstructure of the SAC3595+0.05% Al solder/Cu joint cooled at
0.17.degree. C./s of a selected DSC test. FIG. 3d is an SEM
(scanning electron micrograph) in back scattered electron mode of
the microstructure of FIG. 3c. The Cu substrate is located at the
top and bottom of FIGS. 3a-3d.
[0030] FIG. 4a illustrates a microprobe image of the as-solidified
microstructure of the SAC3595+0.05% Al solder joint and shows a
metastable, intermediate Al-containing rejected solute phase region
adjacent to the Cu.sub.6Sn.sub.5 interfacial layer, while FIG. 4b
illustrates the profile of the Sn, Cu, Ag, and Al concentrations
(wt %) across the joint interface (in .mu.m distance) of FIG. 4a,
5b, 7b (the interface between the Cu substrate and the solder is
where the 100% Cu drops downwardly). In FIG. 4a, 5a, 7a, the Cu
substrate is located on the left side of the microprobe image.
[0031] FIG. 5a illustrates another microprobe image of the
as-solidified microstructure of the SAC3595+0.05% Al solder joint,
while FIG. 5b illustrates the profile of the Sn, Cu, Ag, and Al
concentrations (wt %) across the solder joint of FIG. 5a.
[0032] FIG. 6 illustrates nanoindentation hardness values of the
Cu.sub.6Sn.sub.5 interfacial layer, Cu substrate metal,
Al-containing rejected solute phase region, and the tin matrix
measured within tin dendrites.
[0033] FIG. 7a illustrates a microprobe image of the thermally-aged
(for 1000 hr at 150.degree. C.) microstructure of the SAC3595+0.05%
Al solder joint, while FIG. 7b shows the profile of Sn, Cu, Ag, and
Al content (wt %) across the thermally-aged joint interface (in
.mu.M distance) of FIG. 7a (the interface being where the 100% Cu
drops off).
[0034] FIG. 8 is a bar graph showing undercooling values (.degree.
C.) for SAC 3595 and SAC 3595+0.05 wt % Al with multiple reflow
cycles. For each alloy, five cycles were conducted wherein each
cycle involved raising the temperature from 160.degree. C. to
240.degree. C. with a 30 second dwell followed by cooling from
240.degree. C. to 160.degree. C. at heating and cooling rates of
0.17.degree. C./second.
[0035] FIG. 9 is a bar graph showing Ag.sub.3Sn blades counts per
1000 .mu.m of interface for SAC+Al alloys for each solder alloy
tested with different Al contents.
[0036] FIG. 10 is a scanning electron micrograph (SEM) illustrating
dark, dispersed phase seen on top surface of SAC+0.20Al.
[0037] FIG. 11 is a SEM illustrating scratches and pullout from new
IMC phase indicated by arrows.
[0038] FIG. 12 is a bar graph illustrating prevalence of
Cu.sub.33Al.sub.17 particles per 1000 .mu.m of interface for SAC+Al
alloys that were tested.
[0039] FIG. 13 is a bar graph illustrating nanohardness
measurements showing hardness of Cu.sub.33Al.sub.17 phases and
other solder joint solidification product phases taken in tin and
in Ag.sub.3Sn blade phase regions. Sn matrix(lit.) and
Ag.sub.3Sn(lit.) are published literature values.
[0040] FIG. 14 is a bar graph summary of the undercooling results
for the indicated alloys in one reflow cycle tests in DSC solder
joints.
DESCRIPTION OF THE INVENTION
[0041] The present invention involves reducing the unusually high
undercooling of SAC (Sn--Ag--Cu) solder joints described above,
where there can be difficulty in nucleating Sn solidification as a
pro-eutectic phase, especially during slow cooling, such as
existing for ball grid array (BGA) joints. As mentioned above,
increased undercooling of the solder joints can promote formation
of undesirable pro-eutectic intermetallic phases, specifically
Ag.sub.3Sn "blades," that tend to coarsen radically, leading to
embrittlement of as-solidified solder joints. To this end, the
present invention provides a solder alloy comprising Sn, Ag, Cu,
and Al having an alloy composition controlled to provide a strong,
impact- and thermal aging-resistant solder joint having beneficial
microstructural features described below and substantially devoid
of Ag.sub.3Sn blades. The solder alloy has a relatively low
liquidus temperature and a narrow liquid-solid mushy zone for
solderability.
[0042] In an illustrative embodiment of the invention, the solder
alloy consists essentially of about 3 to about 4 weight % Ag, about
0.7 to about 1.7 weight % Cu, about 0.01 to about 0.25 weight % Al,
and balance consisting essentially of Sn. The solder alloy
preferably exhibits a relatively low solidus temperature of about
217.degree. C..+-.1.degree. C. and narrow liquid-solid mushy zone
with a liquidus temperature not exceeding about 5.degree. C., often
less than 3.degree. C., above the solidus temperature. Other
alloying elements may be present in the solder alloy that do not
substantially affect the melting temperature thereof
[0043] A preferred solder alloy pursuant to the invention consists
essentially of about 3.4 to about 3.6 weight % Ag, about 0.8 to
about 1.1 weight % Cu, about 0.03 to about 0.20 weight % Al, and
balance consisting essentially of Sn.
[0044] A still more preferred solder alloy consists essentially of
about 3.45 to about 3.55 weight % Ag, about 0.9 to about 1.0 weight
% Cu, about 0.04 to about 0.10 weight % Al, and balance consisting
essentially of Sn.
[0045] A still more preferred solder alloy consists essentially of
about 3.45 to about 3.55 weight % Ag, about 0.75 to about 1.0
weight % Cu, about 0.04 to about 0.15 weight % Al, and balance
consisting essentially of Sn.
[0046] Another illustrative embodiment of the invention provides a
Pb-free solder alloy consisting essentially of about 3 to about 4
weight % Ag, 0.95-y weight % Cu, and y weight % Al and balance
consisting essentially of Sn wherein y is about 0.01 to about 0.25
weight %.
[0047] Still another embodiment of the invention provides a still
more preferred solder alloy consists essentially of about 3.45 to
about 3.55 weight % Ag, about 0.80 to about 1.0 weight % Cu, about
0.10 to about 0.20 weight % Al, and balance consisting essentially
of Sn, especially for BGA applications that involve
thermal-mechanical fatigue environments, like avionics.
[0048] The invention also envisions a modification of the alloy
formulation to eliminate the Ag component for situations where
higher solder melting alloys can be tolerated. Such modified solder
alloy embodiments are described below.
[0049] A still further illustrative embodiment of the invention
provides a solder joint and solder process that embody a
Sn--Ag--Cu--Al alloy of the type discussed above wherein the solder
joint has a microstructure that comprises tin dendrites,
interdendritic multi-phase ternary eutectic (between the tin
dendrites), and pro-eutectic Cu.sub.6Sn.sub.5 particles adjacent
and/or within the tin dendrites and that is devoid of Ag.sub.3Sn
blades. This microstructure is achievable at the relatively slow
cooling rates employed for solder paste reflow and BGA solder
processing.
[0050] The as-solidified solder joint microstructure includes an
interfacial layer comprising Cu.sub.6Sn.sub.5 and preferably an
adjacent metastable, intermediate Al-containing rejected solute
region as a zone of intermediate hardness between the hard, brittle
interfacial layer and the softer tin matrix of the solder
microstructure to provide a beneficial hardness gradient
therebetween. The interfacial layer resides between the copper
substrate and the solder of the solder joint.
[0051] The solder joint is formed by the solder being solidified on
an electrical wiring board and/or about copper electrical
conductors in illustrative embodiments of the invention by various
conventional soldering processes including, but not limited to,
solder paste reflow and BGA.
[0052] A thermally-aged solder joint (e.g. aged for 1000 hours at
150.degree. C.) pursuant to the invention has an interfacial layer
thickness that is about the same as the thickness as the
interfacial layer thickness in the as-solidified condition (e.g. no
more than 30% greater in thickness). As a result, the solder joint
is resistant to thermal aging-induced embrittlement.
[0053] For purposes of further illustrating the invention without
limiting it, the present invention is described below with respect
to modifying a near-eutectic alloy, SAC3595 solder alloy (Sn-3.5%
Ag-0.95% Cu, in weight %) as a base by alloying with a fourth
element, Al (aluminum) substituted for part of the Cu to reduce
undercooling of solder joints. In modifying base SAC3595 solder
alloy, Al was alloyed with the base solder alloy to promote
nucleation of pro-eutectic Cu.sub.6Sn.sub.5 within the solder joint
matrix (liquid alloy) in addition to its formation on the substrate
interface, providing additional interfacial area for Sn nucleation.
The Al addition also may strain the lattice of the Cu.sub.6Sn.sub.5
phase, in both pro-eutectic and interfacial layer phases, to make a
more potent epitaxial nucleation catalyst for Sn, thus reducing the
joint undercooling and the potential to form Ag.sub.3Sn blades,
although applicants do not intend or wish to be bound by any theory
in this regard.
[0054] The bulk undercooling measurements for the solder joints
made from the SAC3595+Al alloys that were selected (i.e. Al=0.01%,
0.025% and 0.05% by weight) are summarized in FIG. 1. That is,
solder alloys--Sn-3.5% Ag-0.94% Cu-0.01% Al; Sn-3.5% Ag-0.925%
Cu-0.025% Al, Sn-3.5% Ag-0.90% Cu-0.05% Al in weight % --were
tested. Each alloy was fabricated as a 100-g chill-cast ingot from
component elements of 99.99% purity and drawn into solid wire of
1.7 mm diameter by the Materials Preparation Center of Ames
Laboratory. For each Al concentration level tested, at least seven
repeated trials were used.
[0055] FIG. 1 shows that the Al addition is an active catalytic
addition since these concentrations of Al have relatively lower
undercooling values as compared to the average undercooling of
unmodified SAC3595 base solder alloy. The range of undercooling
values for unmodified SAC3595 is indicated in FIG. 1 by the
left-hand bar at the zero concentration, with the data spread
indicated by the bracket. Al can be seen to have a potent and
consistent nucleation effect. Note that the nucleation temperature
(T.sub.n.) is defined as the onset point of the exothermic
crystallization peak in each DSC thermogram, consistent with the
literature in this field. Also, the bulk undercooling, .DELTA.T, is
defined at the difference between the onset of melting at the
solidus temperature (T.sub.sol) and the onset of nucleation, i.e.,
.DELTA.T=T.sub.sol-T.sub.nuc. Also note that observation on heating
in a differential scanning calorimeter (DSC) of the solidus
temperature for a eutectic alloy is also the singular eutectic
melting point, T.sub.eut, and not just the start of a melting range
between solidus and liquidus temperatures. The DSC apparatus used
was a Pyris 1 power compensating DSC available from Perkin-Elmer
wherein a pre-fluxed copper pan and copper lid accurately simulated
a solder joint in the DSC test. The copper surfaces were cleaned
with methanol and swabbed with flux (Johnson Mfg. No. 1 flux), and
the fluxing action promoted on a hot plate at 180.degree. C. Each
prefluxed pan was loaded with a methonal-cleaned thin disk of the
selected solder alloy that weighed about 15 mg. After mild
crimping, each pan sample was reflowed for one cycle in the DSC
unit by heating at 10.degree. C./min to a peak temperatue of
240.degree. C. for 30 seconds and cooling at 10.degree. C./min
(0.17.degree. C./s) to ambient temperature to simulate BGA reflow
cooling. At least seven separate (repeat) bulk undercooling (AT)
mesaurements (T.sub.eut-T.sub.nuc=.DELTA.T) were performed for each
solder alloy.
[0056] Referring to FIG. 2, the SAC3595+0.01% Al alloy had a
solidus temperature of 216.5.degree. C. and a liquidus temperature
of about 222.degree. C. The SAC3595+0.025% Al alloy had a solidus
temperature of 217.5.degree. C. and a liquidus temperature of
226.5.degree. C. The SAC3595+0.05% Al alloy had a solidus
temperature of 217.degree. C. and a liquidus temperature of
220.degree. C. Note that the SAC3595+0.025% Al alloy appears to
acquire an anomolous higher melting behavior and, as such, is a
less desirable solder alloy choice in this series. However,
subsequent analytical chemistry testing revealed that this
particular alloy did not exhibit the desired composition and that
later experiments showed a very consistent trend in liquidus
temperature and undercooling with neighboring compositions.
[0057] Referring to FIGS. 3a and 3b, the as-solidified
microstructure of the SAC3595+0.01% Al solder joint and the
SAC3595+0.025% Al solder joint, respectively, cooled at
0.17.degree. C./s of a selected DSC test is comprised of tin
dendrites, fine ternary eutectic between the tin dendrites, and
pro-eutectic Cu.sub.6Sn.sub.5 particles adjacent and/or within the
tin dendrites wherein the fine multi-phase ternary eutectic
includes a beta tin matrix with intermetallic phases, such as
Cu.sub.6Sn.sub.5 and Ag.sub.3Sn, distributed in the tin matrix and
wherein the microstructure is devoid of Ag.sub.3Sn blades. An
interfacial layer comprising Cu.sub.6Sn.sub.5 resides between the
copper substrate and the solder in the as-solidified solder joint.
The Cu substrate is located at the top and bottom of the
photomicrograph.
[0058] Referring to FIGS. 3c and 3d, the typical as-solidified
microstructure of the SAC3595+0.05% Al solder/Cu joint cooled at
0.17.degree. C./s of a selected DSC test is comprised of the fine
ternary eutectic between tin dendrites and pro-eutectic
Cu.sub.6Sn.sub.5 particles adjacent and/or within the tin dendrites
and is devoid of Ag.sub.3Sn blades.
[0059] FIG. 4a illustrates a microprobe image of the as-solidified
microstructure of the SAC3595+0.05% Al solder joint and shows an
interfacial layer comprising Cu.sub.6Sn.sub.5 and an adjacent
metastable, intermediate Al-containing rejected solute-rich region
(Sn--Ag--Cu--Al solid solution phase) as a zone between the
Cu.sub.6Sn.sub.5 interfacial layer and the ternary eutectic
microstructure. The Cu.sub.6Sn.sub.5 interfacial layer resides
between the copper substrate and the solder in the as-solidified
solder joint. FIG. 4b illustrates the profile of the Sn, Cu, Ag,
and Al concentrations across the solder joint of FIG. 4a showing
that the rejected solute phase region contains Sn--Cu--Ag--Al.
[0060] FIG. 5a illustrates another microprobe image of the
as-solidified microstructure of the SAC3595+0.05% Al solder joint,
while FIG. 5b illustrates the profile of the Sn, Cu, Ag, and Al
concentrations across the solder joint of FIG. 5a. FIGS. 5a and 5b
confirm that the rejected solute region is present as a zone
between the Cu.sub.6Sn.sub.5 interfacial layer and the ternary
eutectic microstructure. The rejected solute region is adjacent and
separate from Cu.sub.6Sn.sub.5, the interfacial intermetallic layer
in the as-solidified solder joint.
[0061] Referring to FIG. 6, measured nanoindentation hardness
values of the Cu.sub.6Sn.sub.5 interfacial layer, Cu metal
substrate, rejected Al-containing solute region, and the tin matrix
measured within tin dendrites are shown. It is apparent that the
rejected solute region exhibits a hardness intermediate between the
hardness of the hard, brittle Cu.sub.6Sn.sub.5 interfacial layer
and the softer tin matrix. The rejected solute region is about 44%
harder than the tin matrix. The solder joint thus exhibits a
hardness gradient from the hard, brittle interfacial layer toward
the softer tin matrix that improves impact resistance of the solder
joint, consistent with the well known benefits of typical gradient
microstructures in other alloy or composite systems. The
nanoindentation measurments were made using a procedure similar to
that used for microhardness measurement (see J. Mater. Res., Vol.
7, No. 6, June 1992) and used a diamond "cube corner" indent
tip.
[0062] FIG. 7a illustrates a microprobe image of the thermally aged
(for 1000 hr at 15000) microstructure of the SAC3595+0.05% Al
solder joint. In FIG. 7a, the Cu substrate is located on the left
side of the microprobe image. The thermally-aged solder joint has
an interfacial layer thickness that is about the same as the
thickness as the interfacial thickness in the as-solidified
condition (e.g. no more than 30% greater in thickness) so as to
improve thermal aging embrittlement resistance of the solder joint.
The interfacial layer of the thermally aged solder joint comprises
an outer Cu.sub.3Sn layer (adjacent to the Cu substrate) and an
inner Cu.sub.6Sn.sub.5 layer.
[0063] FIG. 7b illustrates the profile of respective Sn, Cu, Ag,
and Al concentrations across the thermally aged solder joint of
FIG. 7a. The aluminum of the rejected solute phase region has
become incorporated into the inner Cu.sub.6Sn.sub.5 interfacial
layer
[0064] The solder alloy pursuant to the invention is useful for
joining electronic assemblies and electrical contacts and to
substitute for Pb-containing solders in all surface mount solder
assembly operations, including solder paste reflow and ball grid
array joints.
[0065] FIG. 8 is a bar graph showing undercooling values (.degree.
C.) for SAC 3595 and SAC 3595+0.05 wt % Al with multiple reflow
cycles. For each alloy, five cycles were conducted wherein each
cycle involved raising the temperature from 160.degree. C. to
240.degree. C. with a 30 second dwell followed by cooling from
240.degree. C. to 160.degree. C. at heating and cooling rates of
1.degree. C./min and 0.17.degree. C./second, respectively. The
average undercooling value for the SAC 3595+0.05 wt % Al was
7.3.degree. C. and generally was much smaller than the undercooling
values for SAC 3595 after multiple reflow cycles.
[0066] Shear strength also was measured by an asymmetric four-point
bend (AFPB) method (see O. Unal, I. E. Anderson, J. L. Harringa, R.
L. Terpstra, B. A. Cook, and J. C. Foley, J. Electron. Mater. 30,
1206 (2001) for larger solder joint specimens made with selected
SAC 3595+Al alloys by hand-soldering with solid solder wire. The
larger specimens (3 mm.times.4 mm.times.75 .mu.m gap) were reflowed
at a peak temperature of 255.degree. C. for 30 seconds and cooled
at 1 .degree. C./s to 3 .degree. C./s to simulate typical surface
mount (paste reflow) soldering processes. AFPB specimens (seven
samples for each condition) were tested as-solidified and after
thermal aging at 150.degree. C. for up to 1,000 h. Microstructural
analysis of the post-AFPB test specimens was performed with SEM on
cross-sectioned metallographic specimens that were polished (ending
with an aqueous slurry of 0.05-.mu.m SiO.sub.2) and ion milled to
provide information on the failure mechanisms.
[0067] Results on solder joints made from Sn-0.95 wt. % Cu (SC95)
were used as a baseline. It should be noted that the large size of
the Cu portion of the shear strength specimens for AFPB testing
prevented them from having their undercooling measured in a
calorimeter. However, the cooling rate for these specimens
(quenched on a massive Cu block) is about nine times faster through
the solidification temperature range, which has a tendency to
promote higher undercooling but does not typically allow time for
massive growth of any Ag.sub.3Sn blades that may nucleate. To cover
this uncertainty in knowledge of the undercooling of these joints,
the microstructure of selected specimens of each alloy was examined
after testing and there was a confirmed absence of Ag.sub.3Sn
blades.
[0068] The most notable feature of the shear strength comparison is
that the as-soldered (unaged) results for SAC3595+0.05% Al and SC95
were nearly identical at about 30 MPa and were lower than SAC3595
(about 41 MPa).
[0069] Further, the thermally aged strength for SAC3595+0.05% Al
was nearly constant at about 30 MPa out to 500 h of aging at
150.degree. C. and only slightly less (29 MPa) at 1000 h. The shear
strength of the of the SAC3595 and SAC3595+0.05% Al alloys seems to
converge at about 30 MPa after 1,000 h of aging at 150.degree. C.
All joints show localized ductile shear failure at about 30 MPa
after 1000 at 150.degree. C.
[0070] A comparison of SC95 and SAC3595+0.05% Al revealed that they
start at about the same moderate shear strength, but the strength
retention for SAC3595 modified by 0.05% Al is significantly better
than the unalloyed SC95. The unaged and aged samples all exhibited
localized ductile shear failures. Inspection of all the shear test
stress-strain curves and microstructural examination of the weakest
post-shear test joints (of seven repeat samples of each alloy)
indicated that the Al additions effectively suppressed the
nucleation and coalescence of pores that can embrittle SAC solder
joints after prolonged high-temperature exposure. The relatively
low initial shear strength and excellent strength retention results
for the SAC 3595+0.05% Al solder appear to relate to the
exceptional stability of the coarse Sn dendrites and fairly stable
interdendritic ternary eutectic microstructure after thermal aging
at 150.degree. C. The Sn--Ag--Cu--Al solder alloy of the present
invention should be useful for low temperature reflow of Pb-free
solder paste and BGA balls (e.g. spheres) as well as other
soldering applications. Another important advantage of the
Sn--Ag--Cu--Al solder pursuant to the invention involves reduction
or avoidance of the formation of Ag.sub.3Sn blades in the
as-solidified solder joint microstructure. Analysis of all of the
solder joint samples for the full range of Al additions revealed
that a minimum of 0.05% by weight Al appears to completely suppress
Ag.sub.3Sn blade phase formation, even at the slow cooling rate
that is common for BGA assembly. This high level of control of the
solder joint microstructure should produce superior results in
board level impact conditions.
[0071] More detailed Ag.sub.3Sn blade counting for the alloys was
conducted on visible Ag.sub.3Sn blades seen protruding from either
the top or bottom of each calorimetric joint interface. Blades of a
length that were .gtoreq.50 .mu.m were recorded in FIG. 9. The
total number of blades was counted per DSC joint, as well as the
total interface length per DSC pan. From these values for
Ag.sub.3Sn blades, the number of blades per 1000 .mu.m of interface
was graphed for each alloy. Ag.sub.3Sn blades were prevalent in the
baseline SAC3595 alloy, as well as for the 0.010% Al concentration.
Blade suppression was seen for the alloys with additions of
aluminum of 0.05 wt % or greater, but blades became visible again
at the higher concentrations. Blades were seen slightly more
frequently on the bottom interface than at the top interface at
lower concentrations, but also had a greater deviation per sample.
At aluminum concentrations of 0.15, 0.20, and 0.25, the Ag.sub.3Sn
blade formation was more consistent per sample and was just as
likely to occur on the top interface as the bottom. The need for an
Ag.sub.3Sn blade suppressant can be seen clearly for the baseline
SAC 3595 alloy where a wide variation in the number of blades can
be seen.
[0072] In addition, as aluminum concentration increased beyond
0.15Al, an increase in ternary eutectic phase fraction was seen, as
well as the appearance of a new small equiaxed (<5 .mu.m) phase.
In the SEM, EDS analysis indicated that the composition of the
small equiaxed particles was slightly enriched in Cu, beyond a 2:1
ratio of Cu:Al. Comparison to the Cu--Al phase diagram and an
extensive analysis of X-ray diffraction results revealed that the
particles were probably Cu.sub.33Al.sub.17 phase. As shown in FIGS.
10 and 11, the particles can be seen either in direct contact with
the Cu.sub.6Sn.sub.5 interfacial layer, attached to the copper
substrate, or adjoined to pro-eutectic Cu.sub.6Sn.sub.5. The
particle phase appears dark in an SEM when observed in
backscattered electron mode. The size of this phase varies from 2-5
.mu.m with an average size near 3 .mu.m.
[0073] WDS analysis of the particles determined the composition to
be 62.2at. % Cu-37.22at. % Al-0.60at. % Sn and it matches closely
with the Cu.sub.33Al.sub.17phase from X-ray results and the initial
EDS analysis. Further examination of the Cu.sub.33Al.sub.17 was
needed, and was conducted on SAC 3595+0.20Al, the alloy that
contained the most Cu.sub.33Al.sub.17 particles of any of the given
alloys (see FIG. 12). When examining Cu substrate/solder interfaces
in the SEM at high magnification on backscattered electron mode, it
can be seen that the new Cu.sub.33Al.sub.17 phase is faceted,
hexagonal in shape and primarily on the top surface (lid side) of
the calorimetric joints. In addition, the phase fraction of
Cu.sub.33Al.sub.17 particles can also be seen in FIG. 12 as a
function of Al content. The total number of particles were counted
per DSC joint, as well as the total interface length per DSC pan.
From these values for Cu.sub.33Al.sub.17 phase particles, the
number of particles per 1000 .mu.m of interface was graphed for
each solder alloy.
[0074] Cu.sub.33Al.sub.17 particles show a different trend with
composition than Ag.sub.3Sn blade formation. In other words, the
phase fraction of particles increases with increasing aluminum
until it reaches an apex at 0.25 wt %, then the Cu.sub.33Al.sub.17
content drops to a level comparable to 0.05Al. Conversely, the
suppression of Ag.sub.3Sn blades is only completely effective for
intermediate levels of Al additions, 0.05Al and 0.10Al. To explain
this behavior partially, one can observe that the formation of
Cu.sub.33Al.sub.17 particles not only depletes the intentional Al
addition, but also reduces the Cu concentration (about 2.times.
faster, in at. %) in the molten solder alloy of the solder joint.
Assuming that the formation of Cu.sub.33Al.sub.17 particles is as
beneficial as the suppression of Ag.sub.3Sn blades, there appears
to be a "sweet spot" in Al content that is centered between about
0.05Al and 0.15Al. However, with higher Al additions, there is
increasingly less available Cu because of the substitutional
alloying approach and because the Cu.sub.33Al.sub.17 particle
formation depletes Cu rapidly. Therefore, less Cu.sub.33Al.sub.17
particles are formed at 0.25Al. It should be noted that a minor
extension of the sweet spot to higher Al could be realized if the
Cu content was not reduced substitutionally with the Al addition,
i.e., maintained at 0.95Cu, without permitting a significant rise
in the solder liquidus temperature.
[0075] A significant piece of data (in FIG. 12) was that 98% of all
particles seen in the SAC3595+Al were found on the top interface
(lid side of the DSC pan sample). Such an observation suggests that
the particles are affected by gravity and that a buoyancy effect
was seen for the particles in the solder joint microstructures.
This buoyancy explanation for segregation of the Cu.sub.33Al.sub.17
particles is consistent with their density, 6.45 g/cm.sup.3, which
is less than the 6.99 g/cm.sup.3 for liquid Sn. Thus, this type of
buoyancy driven segregation of the particles implies that the
Cu.sub.33Al.sub.17 particles nucleate early in the Sn alloy liquid
of the solder joint and float to the top of the joint before
solidification of the Sn and other phases.
[0076] As seen FIG. 11, gouges and unusual scratches were seen in
micrographs of SAC3595+Al alloys. The location of scratches were
often correlated with areas that exhibited pullout of the
Cu.sub.33Al.sub.17 phase. With this idea in mind, nanohardness
measurements were performed on the particles and the other
microstructural phase constituents to determine their relative
hardness (see FIG. 13). This set of nanohardness measurements was
made on a SAC3595+0.20Al DSC joint, using the same cube corner
indenter and load as the previous measurements. The hold time at
maximum load was set at 10 seconds for this set of measurements.
The largest effect of creep during nanoindentation is the initial
penetration, so by increasing the hold time, time dependent effects
like creep are lessened. These measurements of increased accuracy
(shown in FIG. 13) were made in the tin matrix, on an Ag.sub.3Sn
blade, of the Cu.sub.6Sn.sub.5 phase, and on the Cu.sub.33Al.sub.17
particles. Actually, two types of measurements were made for the
Cu.sub.33Al.sub.17 particles; one was for particles engulfed by an
Ag.sub.3Sn blade and the other type were made on particles
supported apparently by the tin solder matrix. The hardness
measurements made on the particles in the tin matrix were
31.4.+-.5.8 GPa compared to particles in the Ag.sub.3Sn blade that
were 49.1.+-.2.5 GPa. The difference in apparent hardness can be
attributed to the hardness of the phase surrounding the particles.
The relatively soft tin matrix has more compliance and is more
ductile than the harder Ag.sub.3Sn IMC. The true hardness of the
particles is most likely closer to the hardness found of
Cu.sub.33Al.sub.17 in Ag.sub.3Sn rather than Cu.sub.33Al.sub.17 in
Sn. Because these Cu.sub.33Al.sub.17 particles do meet the
definition of "superhard" particles, greater than 40-50 GPa, it is
also useful to speculate that they might impart high wear
resistance to a solder joint that is ground to expose the particles
at the top of the joint or that the particles may be extracted and
placed in another matrix for a cutting tool application.
[0077] One worthwhile implication of the observations of hard
particles that float to the top of a solder joint, doped with Al,
is that certain types of thermal-mechanical fatigue (TMF)
environments, especially in BGA joints, could benefit from the
suppression of fatigue crack propagation by this type of
microstructural feature. Hard Cu.sub.33Al.sub.17 particles of the
observed size (about 3 .mu.m) could be very effective at reducing
Sn grain boundary cracking, which is the normal TMF failure
mechanism for Pb-free solder joints, particularly along the top of
BGA joints. Thus, one of the sweet spot alloys, perhaps
SAC3595+0.10Al (or one with higher Cu and slightly higher Al) could
be the optimum alloy for BGA joints that must resist high TMF
conditions.
[0078] Another implication of the work on these alloys is that the
Ag content of the SAC3595+Al alloys does not seem to participate in
the suppression of undercooling by heterogeneous nucleation or in
the generation of the beneficial hard particles. Thus, the
invention envisions to modify the alloy formulation to eliminate
the Ag component for situations where higher melting solder alloys
can be tolerated. For example, an embodiment of such modified
Sn--Cu--Al solder alloy consists essentially of about 0.7 to about
3.5 weight % Cu, about 0.01 to about 0.25 weight % Al, and balance
consisting essentially of Sn. A more preferred solder alloy
consists essentially of about 0.8 to about 3.2 weight % Cu, about
0.03 to about 0.25 weight % Al, and balance consisting essentially
of Sn. A still more preferred embodiment of this solder alloy
consists essentially of about 0.95 to about 3.0 weight % Cu, about
0.15 to about 0.20 weight % Al, and balance consisting essentially
of Sn.
[0079] Still another such modified embodiment involves an alloy
formulation to eliminate the Ag component for situations where
higher solder melting alloys can be tolerated, where another
illustrative embodiment of the invention provides a Sn--Cu--Al
solder alloy consisting essentially of about 3.20-y weight % Cu,
and y weight % Al and balance consisting essentially of Sn wherein
y is about 0.15 to about 0.25 weight %.
[0080] This embodiment of the invention has been tested and FIG. 14
shows the undercooling results in a comparison of the baseline
SAC3595 and the baseline Sn-0.95Cu with one Al level (0.20Al) added
to two different near-eutectic Sn--Cu alloys, along with the
undercooling for SAC3595+0.20Al that was reported earlier. The
suppression of undercooling in these Sn--Cu solder alloys with the
same Al addition is very similar to the SAC+Al alloy and this could
produce the same type of microstructure control benefits in higher
melting (solidus of 227.degree. C.) solder joints.
[0081] While the invention has been described in terms of specific
embodiments thereof, those skilled in the art will appreciate that
modifications and changes can be made thereto within the scope of
the appended claims.
REFERENCES, which are incorporated herein by reference: [0082] K.
N. Tu, A. M. Gusak, and M. Li, J. Appl. Phys. 93, 1335 (2003).
[0083] S. K. Kang, P. A. Lauro, D.-Y. Shih, D. W. Henderson, and K.
J. Puttlitz, IBM J. Res. Dev. 49, 607 (2005). [0084] C. Anderson,
Z. Lai, J. Liu, H. Jiang, and Y. Yu, Mater. Sci. Eng. A 394, 20
(2005). I.E. Anderson, J. Mater Sci. Mater. Electron. 18, 55
(2007). [0085] K.-W. Moon, W. J. Boettinger, U. R. Kattner, F. S.
Biancaniello, and C. A. Handwerker, J. Electron. Mater. 29, 1122
(2000). [0086] D. Swenson, J. Mater. Sci. Mater. Electron. 18,
39-54 (2007). [0087] I. E. Anderson, J. K. Walleser, and J. L.
Harringa, JOM 59, 38 (2007). [0088] W. Liu and N.-C. Lee, JOM 59,
26 (2007). [0089] W. Liu, P. Bachorik, and N.-C. Lee, Proc. 58th
Electron. Comp. Tech. Conf. (IEEE, 2008), pp. 627635. [0090] S. K.
Kang, M. G. Cho, P. Lauro, and D.-Y. Shih, J. Mater. Res. 22, 557
(2007). [0091] I. E. Anderson, J. C. Foley, B. A. Cook, J. L.
Harringa, R. L. Terpstra, and O. Unal, J. Electron. Mater. 30, 1050
(2001). [0092] J. K. Walleser (M. S. thesis, Mat. Sci. & Eng.
Dept., Iowa State University, 2008. [0093] I. E. Anderson and J. L.
Harringa, J. Electron. Mater. 35, 1-13 (2006). [0094] A. Ohno and
T. Motegi, J. Jpn Inst. Metals 37, 777 (1973). Materials
Preparation Center, A. L., US DOE Basic Energy Sciences, Ames,
Iowa, USA. Available from: www.mpc. ameslab.gov. [0095] O. Unal, I.
E. Anderson, J. L. Harringa, R. L. Terpstra, B. A. Cook, and J. C.
Foley, J. Electron. Mater. 30, 1206 (2001). [0096] I. E. Anderson,
J. L. Harringa, and J. K. Walleser, Proc. 4th Int. Braz. Solder
Conf (ASM Int./AWS, 2009), pp. 68-73, ISBN: 978-0-87171-751-1.
[0097] S. K. Kang, W. K. Choi, D.-Y. Shih, D. W. Henderson, T.
Gosselin, A. Sarkhel, C. Goldsmith, and K. J. Puttlitz, JOM 55, 61
(2003). [0098] J.-M. Song, J.-J. Lin, C.-F. Huang, and H.-Y.
Chuang, Mater. Sci. Eng. A 466, 9 (2007). [0099] E. De Monlevade
and W. Peng, J. Electron. Mater. 36, 783 (2007). [0100] K. S. Kim,
S. H. Huh, and K. Suganuma, J. Alloys Comp. 352,226 (2003). [0101]
S. K. Kang, JOM 55, 61 (2003). [0102] I. De Sousa, D. W. Henderson,
L. Patry, S. K. Kang, and D-Y. Shih, Proc. 56th Electron. Comp.
Tech. Conf. (IEEE, 2006), pp. 1454-1461, ISBN: 1-4244-0152-6.
* * * * *
References