The Effect of Heat Input on the Microstructure and Properties

of Nickel Aluminum Bronze Laser Clad with a Consumable

of Composition Cu-9.0Al-4.6Ni-3.9Fe-1.2Mn
C.V. HYATT, K.H. MAGEE and T. BETANCOURT

The effect of heat input in the laser cladding of nickel aluminum bronze was investigated. Nickel
aluminum bronze castings were clad with a consumable of the composition Cu-9.0Al-4.6Ni-3.9Fe-
1.2Mn and exposed to a variety of heat inputs from 42.5 to 595 J/mm. At the lowest heat input, the
deposit microstructure was almost entirely martensitic. Increases in heat input caused the amount of
a to increase. Depending upon heat input, the a was present as grain boundary allotriomorphs,
secondary Widmanstätten a sideplates, and intragranular Widmanstätten a precipitates. The reheated
zones were of lower hardness and, at all heat inputs, consisted of a mixture of grain boundary
allotriomorphs and Widmanstätten a and martensite. Laser cladding improved the corrosion- and
cavitation-erosion resistance of the surfaces but reduced their ductility. The properties of the clad
surfaces depended on heat input.

I. INTRODUCTION Because the intermetallic compounds in nickel aluminum

bronzes are very stable and the a-b phase boundary is near
Laser cladding with nickel aluminum bronze offers the the solidus line, it is very difficult to homogenize these alloys
promise of order-of-magnitude improvements in the sur-
in the solid state.[9,17] Thus, a welding approach is required.
face-sensitive properties of nickel aluminum bronzes and There exists information on the development of micro-
other alloys. It is also a promising repair and reclamation structures in low-heat-input welding of nickel aluminum
method for high-value components. Cladding of nickel alu- bronze.[18–29] Except for some weld simulation studies,[29] all
minum bronze castings with nickel aluminum bronze is of this work has dealt with alloys rich enough in aluminum
particular interest in marine engineering applications. Here, that martensite can be produced following rapid quenching
the development of seawater handling, propulsion, and from the melt. For the alloys examined thus far, in the as-
combat system equipment capable of lasting the life of a deposited material, a martensite with a 9-R structure is
ship would offer enormous cost benefits. Nickel aluminum common in low-heat-input nickel aluminum bronze welds
bronzes are selected for use in high-performance marine over a range of compositions. Other phases, identified by
engineering applications for their ease of fabrication by various workers as Widmanstätten a, bainite, or lath mar-
casting, welding, and forging and for their properties.[1,2] tensite, also form. Weld simulation data also show that, for
Their resistance to corrosion, erosion-corrosion, cavitation-
at least one alloy composition, formation of Widmanstätten
erosion, corrosion-fatigue, stress corrosion cracking, and a can be suppressed with a high enough cooling rate
dealloying are particularly important. So are their toughness Dt8007C–5007C of 0.1 seconds or faster). As well, data show
and strength. These properties impose limitations on the that tempering of martensites in weld-reheated and heat-
design, use, and life of components made of nickel alumi- affected zones probably does not occur at temperatures of
num bronze. For example, the design guidelines for pro- less than roughly 400 7C.[9,30] Transmission electron mi-
pellers made of nickel aluminum bronze are basically croscopy studies, to be published shortly, suggest that,
limited by corrosion fatigue,[2,3,4] the life and inspection in- within the reheated zones between passes, the structure con-
tervals of nickel aluminum bronze valves in seawater han-
sists of a mixture of faulted 9-R–type martensite and coars-
dling systems are limited by dealloying,[5–8] and the damage
ened, light-etching Widmanstätten plates.[18,19,26,27,31] It has
of nickel aluminum bronze impellers and propellers by cav- also been shown that the prior wrought or cast structure
itation-erosion does occur. Improvements in these proper- plays a significant role in the heat-affected zone microstruc-
ties, without degradation of others, could give significant ture. For example, within the heat-affected zone, partial dis-
cost and performance benefits. solution and transformation of prior intermetallics
It has been known for some time that the properties of occurs.[23,24,27,28] A more complete review of microstructural
aluminum bronzes can be modified by solid-state quenching[9] transformations related to low-heat-input welding of nickel
and by rapid solidification and laser surface melting.[10–16]
aluminum bronze is given elsewhere, but a greater under-
standing is required.[31]
Previous work on laser cladding of nickel aluminum
C.V. HYATT, Defence Scientist, is with Defence Research bronzes, especially in heat-input regimes useful in engi-
Establishment Atlantic, Dartmouth, NS, Canada B2Y 3Z7. K.H. MAGEE,
formerly Metallurgical Engineer with The Laser Institute, Edmonton, AB, neering practice, has dealt mostly with microstructural char-
Canada T6E 5J1, is Senior Metallurgical Engineer with CANSPEC Group acterization. Data on properties and on the relationship
Inc., Edmonton, AB, Canada T6P 1N8. T. BETANCOURT, formerly between processing conditions, microstructure, and prop-
Corrosion Scientist with the Nova Scotia Innovation Corp., Dartmouth, erties are lacking. As well, the data that are available have
NS, Canada B2Y 3Z7, is retired.
Manuscript submitted August 29, 1997.
been accumulated on a number of different alloys, using a

REPRODUCED BY PERMISSION
METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1677
OF THE MINISTER OF SUPPLY AND SERVICES CANADA
 Table I. Compositions, Dimensions, and Uses of Coupons

Designation Thickness Specimen Composition

Alloy (UNS) (mm) Cu Al Ni Fe Mn
A C95800 ;25 80.3 8.75 4.8 4.4 1.4
B C95800 ;25 81.1 8.63 4.8 4.0 1.1
C C95800* 6.4 79.3 9 4.3 4.9 0.86
W (ER) 0.25, 0.89 80.8 9.0 4.6 3.9 1.2
(wire) CuNiAl (diameter)
*The material was supplied with this designation. Because iron is greater than nickel, it does not satisfy the chemical analysis requirements.

Table II. Cladding Conditions for Specimens Used in This Work

Laser Travel Wire Heat

Type of Base Power Speed Diameter Defocus Number of Shielding Input
Specimen Alloy (W) (mm/s) (mm) (mm) Layers Gas (J/mm)
C43/1 C 1800 42 0.25 — 1 helium 42.5
C47/1 C 1800 38 0.25 — 1 helium 47.2
C64/2 C 2700 42 0.25 6.4 2 helium 64.3
C64/3 C 2700 42 0.25 6.4 3 helium 64.3
C595/2 C 5000 8.4 0.89 12 2 helium 595
A390/1 A or B 4100 10.6 0.25 12 1 argon 390
A68/1 A or B 2900 42 0.25 7.6 1 argon 68.0
A150/1 A or B 3300 21 0.25 10.2 1 argon 150
A150/2 A or B 3300 21 0.25 10.2 2 argon 150
A150/3 A or B 3300 21 0.25 10.2 3 argon 150

range of different processes and welding geometries. This feeder based on a stepper motor. Precise positioning of the
makes comparisons difficult. This article overcomes the needle tip was possible because it was mounted on a mi-
shortcomings just mentioned by examining the effect of crometer stage. With the fine wire arrangement, shielding
processing conditions in laser cladding on the microstruc- gas was directed at the specimen surface through the pri-
ture and properties of nickel aluminum bronze. One com- mary gas nozzle attached to the laser optics head.
position of commercially available ErCuNiAl weld wire A number of process variables were adjusted by trial and
was used to clad UNS C95000 castings typical of those error to produce the best-appearing deposit. For both wire
used in marine engineering applications. Heat inputs from sizes, wire was fed from the leading edge of the melt at an
42.5 to 595 J/mm were used, and a number of significant angle of 30 deg to the specimen surface. In the 0.89-mm
properties were investigated. wire arrangement, the specimen surface at the start of clad-
ding was located at a position 11.9 mm beyond the focal
point and the wire was aimed at the front of the melt pool.
II. MATERIALS AND METHODS In the 0.25-mm wire arrangement, the specimen surface
A. Castings and Consumables was located at a position 6.4 mm beyond the focal point
and the wire was aimed to intersect the specimen surface
The materials used in this study included cast coupons about 1 mm in front of the melt pool. Information on the
meeting UNS C95800 and ErCuNiAl weld wire. Several processing conditions used to produce each type of speci-
different sizes and compositions of coupons were used. men examined in this work is given in Table II.
These are summarized in Table I. Pretreatment of the cou- More details of the development of the laser welding
pons prior to welding involved machining them flat, glass procedures and their applications to the cladding of com-
bead blasting, washing with a soft bristle brush and soapy plex parts will be given in a future report.[32]
water, and degreasing with acetone.
C. Evaluation Procedures
B. Laser Cladding
1. Potentiodynamic corrosion studies
The cladding experiments were done with a 5 kW CO2 Corrosion behavior was assessed on specimens of as-re-
laser equipped with two different arrangements to feed the ceived alloy A and specimens of alloy A clad with heat
wire to the deposit. One arrangement used 0.89-mm wire inputs of 68 and 390 J/mm, respectively. The three as-re-
and the other arrangement used 0.25-mm wire. The ar- ceived specimens examined were ground to a 600-grit fin-
rangement used to feed and position the 0.89-mm-diameter ish. For each cladding condition, the specimens examined
wire consisted of a gas metal arc welding torch attached to were tested in the as-clad condition. Immediately prior to
a commercial wire feeder. Argon or helium shielding was testing, all specimens were washed in distilled water and
provided through orifices in the torch tip in the normal way. degreased in acetone. The test system was calibrated ac-
In the arrangement used to feed the 0.25-mm-diameter wire, cording to the requirements of ASTM standard G5.[33] Tests
the gas metal arc welding torch was replaced with the tip were done with a three-electrode system using natural sea-
of a hypodermic needle connected to a custom-built wire water at 25 7C which had been purged with nitrogen gas.

1678—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A

Corrosion potentials were determined after equilibrating the entirely clad on one 76.2 3 4.8-mm face. The direction of
test specimens for 1 hour. During the polarization tests, welding was along the length of the specimens. For reasons
voltage was applied continuously from open circuit poten- of material availability, only specimens of type Al50/2 were
tial to 1.1 V at a rate of 0.1 mV/s. investigated. In the tests, one end of the specimen was
clamped in a vice against a bend die and the specimen was
2. Cavitation-erosion experiments bent slowly around the die using stout leather as a backing.
Cavitation-erosion experiments were carried out accord- The test was viewed to be a success if the specimen bent
ing to ASTM standard G-32.[34] The test setup consisted of completely around the bend die without the occurrence of
a piezoelectric vibrator unit equipped with a 12.7-mm-di- cracking detectable by dye-penetrant inspection. The pur-
ameter horn to which a specimen was attached. Duplicate pose of the testing was to compare the ductility of the clad
test specimens were machined from material clad with sin- material to that of the base material. Thus, half the speci-
gle layers at heat inputs of 68 and 390 J/mm and from mens were tested with the clad faces toward the face of the
unclad material. The test surface of each test specimen was die and half were tested with the clad faces away from the
ground to a 600-grit finish, and the specimen was washed face of the die.
in water and then in ethanol in an ultrasonic bath. Follow-
ing air drying, the specimen was weighed to within 50.1 5. Stress corrosion cracking behavior
mg. The test specimen, mounted on the horn of the ultra- Stress corrosion cracking tests were performed using
sonic unit, was immersed in 1.5 L of seawater at 25 5 2 specimens like those described in References 37 and 38 and
7C and oscillated at a frequency of 20 5 0.5 kHz and at shown in Figure 1. These ‘‘tuning-fork specimens’’ al-
an amplitude of 50 mm. Periodically, the experiment was lowed the stresses to be estimated on the nonlaser-clad face
interrupted and the specimen was removed, cleaned, and from strain gage records collected from a dummy specimen.
weighed so that mass loss as a function of time could be Three types of specimens were used. All had been laser
determined. The performance of the test setup was shown clad with two layers on one outer face with a heat input of
to meet the requirements of ASTM standard G-32 by run- 150 J/mm (specimen type Al50/2). One type of specimen
ning a test with annealed nickel 200 in distilled water at 20 had an as–laser-clad finish. A second type was finished by
7C. To ensure consistency, prior to each seawater test, the grinding on metallographic papers to a 600-grit finish. A
seawater was sonicated with a Ni 200 specimen for 30 third type was similarly ground but contained defects in the
minutes to stabilize the gas content of the water. Calcula- region of high stress. These defects were either surface-
tion of the incubation time was obtained by extrapolating breaking porosity, caused by the interaction of the laser-
a tangent from the steepest linear section of the erosion- clad deposit with prior casting defects, or scribe lines.
time curve to the time axis. This intersection with the time Stresses between 60 pct of the yield strength and 140 pct
axis was taken as the incubation time. of the ultimate strength of the base material were induced
by closing the tines of the tuning-fork specimens. Stresses
3. Long-term corrosion behavior were estimated from a relationship between surface strain
To examine long-term corrosion behavior, the specimens and opening of the tines, determined by strain gaging the
listed in Table III were immersed in a tank through which material on the unclad tine. Specimens were immersed in
Halifax Harbor seawater flowed. The specimens were im- flowing Halifax Harbor seawater. Three specimens were re-
mersed for a period of 5 months. Using data from Refer- moved after 1 month and six more were removed after 6
ences 5, 7, and 8, this time meets the requirements of months, the last month of which was spent in stagnant Hal-
ASTM G-4[35] for a valid immersion corrosion test. Speci- ifax Harbor seawater due to a seawater system failure.
mens were 12.4 3 12.4 3 4 mm thick and had a 2.9-mm-
diameter mounting hole drilled through them. The speci-
mens were mounted on a 6-mm-thick sheet of polyacetal III. RESULTS
using nylon screws. Nylon nuts beneath the specimens al-
lowed them to be spaced 1.5 cm away from the board. The A. Metallography
head of the screw and the nylon nut beneath the specimen
created small crevices on the specimen. Crevices were de- Low-magnification optical micrographs of specimens
liberately included on some specimens by putting a poly- prepared by cladding with multiple layers are shown in Fig-
acetal plate the size of the specimen over the surface of the ure 2, while low-magnification optical micrographs of spec-
specimen. Previous researchers have shown that coupling imens prepared by cladding with single layers are shown
nickel aluminum bronze to titanium alloys produces one of in Figure 3. In both cases, microhardnesses measured on
the worst galvanic conditions seen in practice.[5,7] Thus, a the specimens are also shown graphically. In this article,
number of specimens were screwed to a block of Ti-6Al- only the structures of the as-deposited and reheated weld
4V with dimensions of 12.5 3 37.5 3 4 mm. The area of metal are discussed. The reasons for this simplification are
the Ti alloy block was more than 6 times the area of the that the structure of the base casting has a large effect on
specimens. In most cases, the clad faces (only one face was the nature of the heat-affected zone and that, in the exper-
clad) faced up. The exceptions were a few cases where a iments on properties, it is the properties of the deposit, not
clad face was coupled with its face toward a Ti alloy block. those of the base metal heat-affected zone, which are being
investigated.
4. Bend tests Before proceeding to discuss the results, it is necessary
The ductility of a typical laser-clad specimen was as- to explain the difference in the thickness of deposit on the
sessed by performing bend tests according to ASTM stan- A and C series alloys and to comment on the significance
dard E 290-87.[36] Specimens for bend testing had of this to the development of microstructure. The C series
dimensions of 76.2 3 4.8 3 1.56 mm. The specimens were specimens were clad using a helium gas shield, while the

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1679

 Table III. Specimens Used in Dealloying Experiments

Metallographic
Specimen Coupling Preimmersion Postimmersion Weight Loss Observations—
Origin and Crevices Weight (g) Weight (g) (g) Depth of Pitting
C(UT) none 3.6851 3.6839 0.0012 20 mm
C64/2 none 4.1290 4.1262 0.0028 most .15 mm, one
60 mm
A150/2 none 3.5135 3.5111 0.0024 15 mm
A390/1 none 4.0699 4.0677 0.0022 20 mm
A(UT) none 3.6830 3.6807 0.0023 100 mm
A(UT) none 3.7468 3.7426 0.0042 40 mm
A(UT) none 3.7834 3.7805 0.0029 30 mm
A(UT) to Ti 4.2973 4.2936 0.0037 100 mm
A(UT) to Ti 3.7861 3.7802 0.0059 100 mm
A150/2 to Ti 3.14842 3.4814 0.0028 15 mm
A150/2 coupled clad face 3.3646 3.3598 0.0048 10 mm
down to Ti
A(UT) nylon block 4.2736 4.2719 0.0017 20 mm
covering upper
surface
A150/2 nylon block 3.5866 3.5849 0.0017 10 mm
covering clad
surface

A series specimens were clad using an argon gas shield.

For reasons to be discussed in another article, this allowed
larger amounts of material to be deposited on the C series
specimens than on the A series specimens.[32] For the pur-
poses of this article, it is significant that differences in mi-
crostructures and microhardnesses in specimens, resulting
from different amounts of dilution in specimens produced
by cladding with either different deposition rates or differ-
ent numbers of passes, appear to be minimal. This is based
on the occurrence of similar hardnesses and microstructures
in single and multipass welds and in welds produced at the
same heat input but with different deposition rates. This
probably occurs because the composition of the castings is
similar to that of the weld deposit, and minimal dilution is
expected. More significant dilution effects are observed
when dissimilar substrates are clad.[32] Fig. 1—Tuning fork-type stress corrosion specimen.
Significant changes in hardness occurred with increases
in heat input. As shown in Figure 4, peak deposit hardness
increased to a maximum Vickers hardness (HV) of 540 at products as well as the products of autotempering.[31] Thus,
a heat input of about 64 J/mm. The peak deposit hardness for this article, we will designate the dark etching phase as
dropped after this. In contrast, as the graphs in Figures 2 the martensitic phase, realizing that the actual structures
and 3 show, the lowest hardness in the reheated zones was may be more complex in some cases. It is significant to
between 200 and 260 HV regardless of heat input. note that we are not the first to encounter this difficulty.
Optical micrographs of the as-deposited material are Workers who deal with solid-state transformations in nickel
shown in Figure 5. Those of the corresponding reheated aluminum bronzes have often used the terminology b* to
zones are shown in Figure 6. Etching with 5 g FeCl3 and describe all metastable b transformation products from
15 mL of HCl3 dissolved in 60 mL of ethanol revealed a quenching. Because the transformation products being con-
darker etching phase and a lighter etching phase. The sidered here may be different than the solid-state ones, we
lighter etching phase occurred as grain boundary allotrio- have chosen not to adopt this convention.
morphs and also, sometimes, throughout grains, with a The as-deposited material changed with heat input. At
Widmanstätten morphology. The darker etching phase be- the lowest heat input investigated, 42.5 J/mm (Figure 5(a)),
came progressively easier to etch on specimens produced most of the as-deposited material seemed to be martensite.
with increasing heat input and showed more internal detail. Only a small amount of grain boundary a formed. At a
We believe the white etching phase to be the a phase solid heat input of about 47.2 J/mm (Figure 5(b)), more than half
solution, which is stable in these alloys at low temperatures, the prior b grain boundaries had grain boundary allotriom-
and in this article we will call it the a phase. Designating orphs of a on them. As well, secondary Widmanstätten
the dark etching phase is more difficult. It certainly contains sideplates formed on many of the prior b grain boundaries
some martensite. However, it has been suggested that, at upon which grain boundary allotriomorphs of a had
least at a few heat inputs, it also contains some bainitic formed. Deposits produced with a heat input of 64.3 J/mm

1680—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A

 (a) (b)

(c)
Fig. 2—Optical micrographs and microhardness scans on specimens clad with multiple passes: (a) specimen of type C64/3 (64.3 J/mm); (b) specimen of
type A150/2 (150 J/mm); and (c) specimen of type C595/2 (595 J/mm).

(Figure 5(c)) had grain boundary allotriomorphs of a on may have occurred, although a darker-etching martensitic
most prior b grain boundaries. As with the 47.2 J/mm case, phase is still present. A deposit produced with a heat input
secondary Widmanstätten a sideplates formed on many of of 390 J/mm was similar, although the volume fraction of
the prior b grain boundaries upon which grain boundary a a seems to have increased, primarily by the formation of
had formed. Deposits produced at a heat input of 150 J/mm what appears to be intragranular Widmanstätten plates and
(Figure 5(d)) had all prior b grain boundaries outlined with increased secondary Widmanstätten plate formation. The
a. Secondary Widmanstätten sideplates were pronounced, highest heat input investigated in this work, 595 J/mm, pro-
and some intragranular Widmanstätten precipitation of a duced a microstructure with a similar volume fraction of

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1681

 (a) (b)

Fig. 3—Optical micrographs and microhardness scans on specimens clad

with single passes: (a) specimen of type C47/1 (47 J/mm); (b) specimen
of type A68/1 (68 J/mm), and (c) specimen of type A 390/1 (390 J/mm).
Note the specimens shown in Figs. 3(b) and (c) had been subjected to
cavitation erosion prior to sectioning.

(c)

1682—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A

 current, which are proportional to the corrosion rate, are
given in Table IV. These results suggest that specimens
clad with a heat input of 68 J/mm will corrode about 10
times more slowly than the as-received material.

C. Cavitation-Erosion Resistance
The results of the cavitation-erosion experiments are
shown in Figure 8. This figure shows the mean depth of
erosion for laser-clad material produced with a low heat
input (68 J/mm), clad material produced with a high heat
input (390 J/mm), and that for the base material. The max-
imum cavitation-erosion rate and the incubation times for
the three types of specimens tested are shown in Table V.
The results reveal that the resistance to cavitation-erosion
(based on maximum erosion rate) was increased by a factor
of 5 by laser cladding with a heat input of 68 J/mm and by
a factor of 2.5 by laser cladding with a heat input of 390
J/mm. They also show that low-heat-input cladding with a
heat input of 68 J/mm doubles the incubation time, while
high-heat-input cladding has no effect on incubation time
relative to that of the bulk material.
Fig. 4—Peak deposit hardness as a function of heat input.
Sectioning of the cavitation-erosion specimens showed
that attack was fairly extensive on the base alloy (Figure
martensite to the 390 J/mm case. The 595 J/mm heat input 9(a)), but was concentrated at the heat-affected zones of the
did seem to produce more intragranular Widmanstätten clad specimens (Figures 9(b) and (c)). The higher-heat-in-
plates. put specimen (A390/1—Figure 9(b)) was affected more se-
Compared to the as-deposited material, and except for verely than the lower-heat-input specimen (A68/1—Figure
the size of the zone, the reheated zone seemed to change 9(c)). This may be primarily due to hardness effects, but
less with heat input. This might have been expected from might also be affected by a change in the nature and density
the relatively similar minimum reheated zone hardnesses. of Widmanstätten plate boundaries. A similar effect has
In this article, we consider only the coarsest part of the been reported to affect the toughness of microstructurally
reheated zone. The reheated zones corresponding to the as- similar Ti alloys.[39] The presence of crack-stopping plate-
deposited materials just discussed are shown in Figure 6. type boundaries in the clad specimens may explain why
In the reheated zones, for all heat inputs, the prior b grain cracks running perpendicular to the surface occurred in the
boundaries are completely covered with grain boundary al- unclad material (refer to the arrow in Figure 9(a)) but did
lotriomorphic a. Also, in all cases, there is some secondary not occur in the clad specimen. It is also interesting to note
Widmanstätten plate formation and extensive formation of that, on the low-heat-input specimen, cavitation damage
what could be described as a basketweave pattern of (prob- seems to have occurred more deeply at and around prior b
ably intragranular) Widmanstätten a.[29] From Figure 6, it boundaries, most of which were covered with grain bound-
is apparent that the main change in the reheated zone with ary allotriomorphic a (refer to the arrow in Figure 9(c)).
increasing heat input is an increase in coarseness. It is also
interesting that, at the two highest heat inputs examined
(390 J/mm, Figure 6(e), and 595 J/mm, Figure 6(f)), a D. Long-Term Corrosion Behavior
clearly equiaxed two-phase a plus b zone has developed in
the reheated zone adjacent to the fusion line. The long-term corrosion specimens were coated with ma-
rine growth after removal from the immersion tanks. This
B. Potentiodynamic Corrosion Behavior heavy coating of marine growth is thought to have pro-
duced a fairly severe seawater corrosion environment. The
The results of the corrosion experiments are summarized results of these experiments, presented in terms of weight
in Table IV and Figure 7. As shown in Table IV, the cor- loss and maximum depth of pitting or preferential attack,
rosion potential of each specimen type was similar and typ- are summarized in Table III. Weight loss data provide lim-
ical of most copper alloys.[2] Grinding the surface of the ited insight because only one face of the specimens was
clad specimens to a 600-grit finish did not significantly clad. Except for very local effects less than 10 mm deep on
change the corrosion potential from that of the as-clad con- the treated and untreated alloy C, there was little evidence
dition. Polarization curves are shown in Figure 7. Pitting, of dealloying, except on a few specimens which had de-
as suggested by the instability in current on one of the veloped a coppery surface hue. Thus, discussion of deal-
polarization curves for the 390 J/mm–clad specimens (Fig- loying effects will need to be deferred until specimens
ure 7(b)) and confirmed by examination of this specimen, immersed for a longer term can be examined.
did occur on one of the 390 J/mm specimens. Grain bound- The clad specimens which were exposed in the uncoup-
ary attack occurred on the as-received material. The cor- led condition without a nylon crevice all exhibited minor
rosion current was estimated by extrapolating the anodic pitting in the as-deposited material. Similar results were
curve to the open circuit potential. Values of the corrosion observed on specimens exposed coupled to Ti. A severe

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1683

Fig. 5—Optical micrographs of as-deposited material: (a) specimen C43/1 (heat input of 42.5 J/mm); (b) specimen C47/1 (heat input of 47 J/mm); (c)
specimen C64/3 (heat input of 64.3 J/mm); (d ) specimen A150/2 (heat input of 150 J/mm); (e) specimen A390/1 (heat input of 390 J/mm) and ( f )
specimen C595/1 (heat input of 595 J/mm).

example is shown in Figure 10(a). In all cases, the darker- icant change in weight loss with heat input, and the mea-
etching martensite was attacked preferentially, often leaving sured weight losses of the clad materials were consistent
the lighter-etching grain boundary and Widmanstätten a with those of the base materials. However, two changes in
phases standing proud on the surface. There was no signif- location of pitting did occur with increases in heat input.

1684—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 6—Optical micrographs of reheated zone of deposit: (a) specimen C43/1 (heat input of 42.5 J/mm); (b) specimen C47/1 (heat input of 47 J/mm); (c)
specimen C64/3 (heat input of 64.3 J/mm); (d ) specimen A150/2 (heat input of 150 J/mm); (e) specimen A390/1 (heat input of 390 J/mm); and ( f )
specimen C595/2 (heat input of 595 J/mm).

First, there seemed to be an increased tendency toward pit- C64/2. This pitting is shown in cross section in Figure
ting under the crevice in the lowest-heat-input specimen. 10(b). Its depth is approximately 0.1 mm. It is interesting
The worst pitting seen in this work was observed in this to note that similar pitting was also seen under the crevice
region under the screw head on the clad face of specimen of a specimen of surface-melted (i.e., autogenously welded)

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1685

 Table IV. Corrosion Potentials and Currents A390/1 and comparable to the one severe pit found on alloy
C64/2. Adding a nylon cap worsened the amount of pitting
Corrosion Potential vs
SCE at 23 7C (mV)
only slightly on both specimen A150/2 and the untreated
Heat Corrosion alloy A. This result can be rationalized on the basis that
Input Individual Current even the uncoupled specimens had similar crevices created
Specimen (J/mm) Test Average (mA/cm2) by the presence of marine growth. It was in the case of
Alloy A 1 NA* 2346 — 0.46 titanium alloy coupling that the worst results were ob-
Alloy A 2 NA 2337 — 0.4 served. Titanium-coupled alloy A had extensive 100-mm-
Alloy A 3 NA 2332 2338 5 8 0.25 deep pits and was the most severely attacked type of spec-
A68/1 1 68 2362 — 0.014 imen in this work, as shown in Figure 10(f), while, as dis-
A68/1 2 68 2345 2353 5 9 0.048
A390/1 1 390 2342 — pitting
cussed in the previous paragraph, titanium-coupled
A390/1 2 390 2361 2351 5 10 0.16 specimens of type A150/2 behaved only slightly worse than
the equivalent uncoupled specimens, whether they were ex-
*NA—not applicable.
posed clad face toward or away from the titanium block.

UNS C95800 exposed together with the clad specimens dis- E. Bend Tests
cussed in this article. This specimen had been produced The bend tests showed that the clad material was much
with a heat input of 20 J/mm.[21] Its composition was Cu- less ductile than the unclad material. When tests were con-
9.1Al-4.4Ni-3.8Fe-1.2Mn, nearly identical to that of the ducted so that the clad surface was in tension, cracking did
welding consumable used in this work. These results sug- not occur at a bend radius of 28.7 mm but did occur at a
gest that surfaces produced with very low heat inputs are radius of 25.4 mm. When tests were conducted so that the
more susceptible to crevice effects, although more work is unclad surface was in tension, cracking did not occur at a
required to understand these observations. It is also possible bend radius of 9.6 mm but did occur at a bend radius of
that these specimens exhibited a greater tendency toward 6.4 mm. There were also differences in how cracking oc-
crevice corrosion because all welding oxides were not re- curred on the two faces. Cracks on the clad specimens were
moved prior to immersion. The second change in location audible and traveled across the whole clad face, while
of pitting with heat input was that only in the highest-heat- cracks on unclad specimens were not audible and were less
input specimen (390 J/mm—specimen A390/1) tested did extensive. The results show that the clad surface has a
pitting occur at the reheated zone in addition to in the de- lower toughness than the base material.
posit. This is shown in Figure 10(c). The presence of an
equiaxed a plus martensite region in the reheated zone just
adjacent to the fusion zone of the higher-heat-input speci- F. Stress Corrosion Cracking Resistance
mens does not seem to have played a role in causing pitting Three stress corrosion cracking specimens were removed
to occur here. from the seawater immersion tank after 1 month. Despite
On all specimens, the interface between the clad deposit stresses in excess of the yield strength of the base material,
and the base material was exposed to seawater, although in no evidence of stress corrosion cracking was noted on ei-
a number of cases, marine growth created a crevice on top ther the treated or the untreated tine. In none of the nine
of this area. In none of the specimens examined was there stress corrosion specimens examined was evidence of
significant preferential attack of this region. An example of cracking found by visual or die-penetrant inspection. Three
the appearance of this region after testing is shown in Fig- of the specimens had developed severe pitting. Each of the
ure 10(d). pitted specimens had been exposed to stagnant Halifax
A laser-clad specimen of type A150/2, immersed with Harbor seawater. Pitting occurred on both the clad and un-
the clad surface covered by a nylon pad, had a slightly more clad faces, usually under a green deposit. This deposit,
coppery surface appearance but was otherwise similar to when dry, contained significant amounts of copper and
the specimen of type A150/2 which was exposed without chlorine, lesser amounts (5 to 15 pct by weight) of iron and
a pad. A specimen of type A150/2 immersed bolted to a aluminum, and trace amounts (less than 1 pct by weight)
titanium block, with the clad face up, exhibited similar, of sulfur and calcium. Powder X-ray diffraction showed
although slightly more, pitting than an uncoupled specimen. that the copper hydroxychlorides Cu(OH)Cl, Cu3(OH)2Cl4,
A specimen of type A150/2 immersed bolted with its clad Cu3(OH)2Cl4z2H2O, and Cu11(OH)14Cl8z6H2O were present
face against a titanium block also exhibited more pitting. in the deposit. A number of different surface finishes were
However, in this case, pitting occurred at the reheated zone investigated in this work. It is significant that pitting did
as well. not develop on areas of the clad surface which had been
In general, the behavior of the clad materials in the long- ground, but only on regions which had the high-temperature
term exposure tests are superior or similar to that of the oxide on the surface intact.
base alloy A, the only base alloy for which a large number
of specimens were available. Alloy A was prepared as a
15.2-cm-diameter by 30.5-cm-long sand casting and, thus, IV. DISCUSSION
had a coarse structure. When exposed uncoupled and with-
out a nylon cap, alloy A exhibited maximum pitting, or The results of this work show that, except for a signifi-
preferential corrosion of between 40 and 100 mm. An ex- cant reduction in ductility, laser cladding either improved
ample is shown in Figure 10(e). These are higher than the or did not measurably degrade the properties assessed in
values observed for the specimens of type A150/2 and this work. It also showed that microstructure was strongly

1686—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A

 (a) (b)

Fig. 7—Corrosion behavior of nickel aluminum bronze: (a) as-received

alloy A; (b) specimen of type A68/1 (heat input of 68 J/mm); and (c)
specimen of type A390/1 (heat input of 390 J/mm).

(c)

Table V. Summary of Cavitation-Erosion Test Results

Heat Input Maximum Erosion Incubation

Specimen (J/mm) Rate (mm/h) Time (Min)
Alloy A NA 7.6 69
A68/1 68 1.5 114
A390/1 390 3.0 70

of a peak deposit hardness at a heat input of 62.5 J/mm is the

result of precipitation hardening. The order of cavitation-ero-
sion resistance observed in this work was laser-clad at 68
J/mm . laser-clad at 390 J/mm . base material.
This is the same order as the average material hardness,
although the differences in ductility of the base and laser-
clad material did not make the observed order of cavitation
resistance predictable. It is intriguing that the base material
Fig. 8—Cavitation-erosion behavior of as-received alloy A (labeled base and the higher-heat-input clad material have the same in-
material) and clad specimens of types A390/1 (labeled high heat input) cubation time, while the higher-heat-input clad material is
and A68/1 (labeled low heat input). more resistant to cavitation-erosion. In contrast, the low-
heat-input case had a longer incubation time than the base
dependent on heat input. This was especially true in the as- material and an improved resistance to cavitation-erosion.
deposited region. The reheated zone microstructure and It has long been noted that there is a relationship between
hardness seemed to be less sensitive to heat input. corrosion-fatigue limit and incubation time,[2] so, at least
The different microstructures produced by different heat in- for the low-heat-input case, an increase in the corrosion-
puts had different properties. We surmise that the observation fatigue limit might be expected. The observed localization

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1687

 interaction between corrosion-fatigue and cavitation-ero-
sion. For both clad conditions subjected to cavitation-ero-
sion testing, an improvement in corrosion fatigue life in the
so-called mortal zone is also expected, based on published
relationships between cavitation-erosion data and mechan-
ical properties.[40]
The low-heat-input deposit microstructure, which fa-
vored cavitation resistance, seemed to be more prone to
pitting in the long-term corrosion experiments than the
higher-heat-input deposits. It is interesting that only at
higher heat inputs did pitting of heat-affected zones occur.
It is also interesting that only in the higher-heat-input case
did the potentiodynamic corrosion studies suggest a ten-
dency to pitting. The presence of oxide from welding
seemed to be associated with much, and perhaps all, of the
pitting observed. Explanation of these phenomena requires
more detailed metallographic and electrochemical studies,
but is an important part of optimizing deposit processing
and composition to give optimum properties.
The results in this article are for a commercially available
consumable composition. While improvements in proper-
ties over the bulk occur over a wide range of heat inputs,
both structure and properties were highly dependent upon
heat input. Comparing the results of this study to work with
a consumable of composition Cu-10.6Al-5.1Ni-1.3Mn
shows that both the peak hardness and the heat input at
which the peak hardness is achieved are functions of heat
input.[41] Thus, to reduce the sensitivity to heat input, and
to make greater property improvements, the effect of the
consumable composition needs to be examined and the con-
sumable composition needs to be optimized. In terms of
microstructure, the goals of such an approach should be to
make the as-deposited material microstructure less sensitive
to heat input and to make the microstructure and hardness
of the reheated material similar to that of the as-deposited
material. In terms of properties, the goals should be to in-
crease ductility and avoid pitting sensitivity in either the as-
deposited or the reheated material, while retaining or
improving upon increases in resistance to cavitation-erosion
and general corrosion.

V. CONCLUSIONS
The microstructures of both the as-deposited and re-
heated material in a laser-clad deposit produced using a
welding wire of composition Cu-9.0Al-4.6Ni-3.9Fe-1.2Mn
were sensitive to heat input. At the lowest heat input in-
vestigated (42.5 J/mm), the as-deposited material had a pre-
dominately martensitic microstructure, although a few prior
b boundaries were outlined with grain boundary allotriom-
orphic a. As heat input increased, the fraction of prior b
grain boundaries which were outlined with a increased. At
a heat input of 47.2 J/mm, secondary Widmanstätten a be-
gan to form. The amount of secondary Widmanstätten a
Fig. 9—Optical micrographs of cavitation erosion specimens: (a) as-
received alloy A; (b) specimen of type A390/1 (labeled high heat input
precipitation increased with heat input. By a heat input of
in Fig. 8); and (c) specimen of type A68/1 (labeled low heat input in Fig. 150 J/mm, primary intragranular Widmanstätten a had be-
8). gun to form in the as-deposited material as well. The
amount and coarseness of Widmanstätten a precipitation
increased with heat input. The hardness of the as-deposited
of cavitation-erosion attack at the reheated zones of the la- material reached a maximum at a heat input of 64 J/mm,
ser-clad specimens may affect the validity of the correlation probably because of a secondary precipitation-hardening ef-
just mentioned and may have implications for in-service fect.

1688—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 10—Micrographs of long-term corrosion specimens: (a) specimen of type A150/2 exposed coupled to Ti; (b) specimen of type C64/2, in region under
nylon screw head; (c) specimen of type A390/1, showing small pit in heat-affected zone; (d ) surface where deposit meets base material on a specimen of
type A150/2 exposed coupled to Ti; (e) as-received alloy A exposed uncoupled; and ( f ) as-received alloy A exposed coupled to Ti-6A1-4V.

Over the range of heat inputs investigated in this work, reheated zones was always below HV 300, regardless of
the reheated zone always contained Widmanstätten and heat input.
grain boundary a. As well, the minimum hardness of the The properties of the deposit also depended strongly on

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 29A, JUNE 1998—1689

heat input. Cavitation-erosion resistance was higher in a Materials Technology, J.R. Matthews, ed., Defence Research
Establishment Atlantic, Dartmouth, NS, Canada, pp. 226-52.
specimen produced at lower heat input, and the location 22. C.V. Hyatt and K.H. Magee: Proc. Advanced Methods of Joining New
where pitting occurred seems to depend on heat input. Of Materials II, The American Welding Society, Miami, FL, 1994, pp.
the properties investigated, only ductility was adversely af- 111-26.
fected by laser cladding. Most properties were improved. 23. D.E. Bell: Master’s Thesis, The Pennsylvania State University,
University Park, PA, 1993.
24. K. Petrolonis: Master’s Thesis, The Pennsylvania State University,
University Park, PA, 1993.
ACKNOWLEDGMENTS 25. P.J. Oakley and N. Bailey: Proc. Int. Conf. on Power Beam
Technology, The Welding Institute, Cambridge, United Kingdom,
The technical assistance of Messrs. Bob Armstrong and 1986, pp. 301-14.
Bob Johnson is gratefully acknowledged. 26. F. Hansan, G.W. Lorimar, and N. Ridley: Proc. Conf. on Phase
Transformations, Cambridge, United Kingdom, 1987, pp. 131-34.
27. D.E. Bell, K. Petrolonis, and P.R. Howell: Proc. Int. Trends in
Welding Science and Technology, S.A. David and J.M. Vitek, eds.,
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1690—VOLUME 29A, JUNE 1998 METALLURGICAL AND MATERIALS TRANSACTIONS A