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Introduction
Ultrasonic additive manufacturing(UAM) is a solid-state joining
process inwhich thin metallic tapes are ultrasonicallywelded on top
of one another and periodi-cally machined to create a final part.
Aschematic illustration of the process isshown in Fig. 1. Along
with progressivebuilding of a block through seam welding(Fig. 1A),
a milling process (Fig. 1B) is usedas required to create holes or
channels be-fore welding the subsequent layers. Themachining
operation is also used periodi-cally to produce a flat surface to
ensureproper dimensions of the finished build.This process offers
many benefits over tra-
ditional fusion welding processes such as al-lowing for complex
shapes and designs,having a significantly lower processing
tem-perature, allowing for embedded materialsand channels, and
offering the capability ofjoining dissimilar materials that are
other-wise difficult or impossible due to UAMbeing a solid-state
process.
The majority of research on UAM is cur-rently focused on
optimizing processing pa-rameters (Refs. 16) and characterizing
thequality and microstructure (Refs. 79) ofthe resulting builds.
The four main param-eters in UAM are sonotrode amplitude,
travel speed, normal force applied, and pre-heat temperature.
Increasing the ampli-tude, normal force, and preheat tempera-ture,
while decreasing the travel speed,generally increases the quality
of the bonds.However, above a threshold for each pa-rameter, no
further gains are realized (Ref.5). The threshold effects with
respect tosonotrode amplitude and normal force aremost likely due
to the machine not being ca-pable of delivering enough power to
sustainthe ultrasonic vibrations at the higher am-plitudes and
forces. Additional gains inbond quality may be possible with a
higher-power system, allowing for higher ampli-tudes of vibration
and forces. For most cur-rent UAM machines, optimum parametersare
approximately 1821 m amplitude,2550 mm/s travel speed, preheat
of65150C, and normal forces between 800and 1500 N. Peel tests
(Refs. 1, 2, 4, 6), fiberpush- out testing (Ref. 3), and
microhard-ness and nanohardness tests (Refs. 2, 9, 10)have been
conducted to further the under-standing of this additive
manufacturingprocess. These tests are often done alongwith
parameter development to comparethe bond quality between different
builds.Voids are often present in UAM builds andcan be quantified
with linear weld density(LWD). The LWD is defined as the lengthof a
particular interface that appears prop-erly bonded divided by the
total interfacelength inspected. The LWD is often used asa test to
determine optimum processing pa-rameters (Refs. 1, 2, 4, 5, 11). It
is generallyagreed that to improve the bond quality ofUAM builds,
LWD must be kept as high aspossible. In most UAM builds, LWD
den-sity ranges from 40 to 95%. Ram et al. (Ref.5) and Johnson
(Ref. 7) theorized voidsform in UAM builds due to the
sonotrodetransferring its texture to the workpiece.This results in
a situation where the top ofeach interface is smooth, but the
bottom is
Microstructural Characterization of BondingInterfaces in
Aluminum 3003 Blocks
Fabricated by Ultrasonic AdditiveManufacturing
A look at linking microstructure and linear weld density to the
mechanicalproperties of ultrasonic additive manufacturing builds as
well as analyzing their
properties with different microscopy and testing methods
BY D. E. SCHICK, R. M. HAHNLEN, R. DEHOFF, P. COLLINS, S. S.
BABU, M. J. DAPINO, AND J. C. LIPPOLD
KEYWORDS
Ultrasonic AdditiveManufacturing (UAM)
Linear Weld Density (LWD)Scanning Electron Microscopy(SEM)
Shear StrengthAl 3003-H18Transmission Electron
Microscopy
D. E. SCHICK, S. S. BABU ([email protected]),and J. C. LIPPOLD are
with the Department ofMaterials Science and Engineering Welding
En-gineering Program, and R. M. HAHNLEN and M.J. DAPINO are with
the Department of Mechani-cal Engineering at The Ohio State
University,Columbus, Ohio. R. DEHOFF, formerly at the De-partment
of Materials Science and Engineering, iscurrently at Oak Ridge
National Laboratory, OakRidge, Tenn. P. COLLINS, formerly at the
Depart-ment of Materials Science and Engineering, is cur-rently at
Quad Cities Manufacturing Laboratory,Rock Island, Ill.
Ultrasonic additive manufacturing (UAM) is a process by which
hybrid and near-net-shaped products can be manufactured from thin
metallic tapes. One of the main con-cerns of UAM is the development
of anisotropic mechanical properties. In this work,the
microstructures in the bond regions are characterized with optical
and electron mi-croscopy. Recrystallization and grain growth across
the interface are proposed as amechanism for the bond formation.
The presence of voids or unbonded areas, which re-duce the
load-bearing cross section and create a stress intensity factor, is
attributed tothe transfer of the sonotrode texture to the new foil
layer. This results in large peaks andvalleys that are not filled
in during processing. Tensile testing revealed the weld inter-face
strength was 15% of the bulk foil. Shear tests of the weld
interfaces showed almost50% of the bulk shear strength of the
material. Finally, optical microscopy of the frac-ture surfaces
from the tensile tests revealed 34% of the interface area was
unbonded.
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rough. To combat this, it has been demon-strated that milling
between layer depositsto provide a smooth-to-smooth interfacecan
eliminate voids, achieving 98% LWD(Ref. 5). However, no tensile,
peel, or otherquantitative measurement of bond qualitywas done to
verify the bond quality.
Several researchers have made prelim-inary attempts at
mechanical and finite el-ement modeling (FEM) of UAM weld-ments.
Doumanidis and Gao (Ref. 10)used an analytical model combined
withexperimental data to produce an FEM of
UAM useful in simulating different mate-rial combinations,
embedding of materi-als, and the production of complex parts.This
model also proved useful in deter-mining ideal geometry for the
sonotrodeand other components. Zhang et al. (Ref.12) developed a
three-dimensional FEMfor ultrasonic spot welding that evaluatedthe
ever-changing parameters at eachnode including normal stress, heat
gener-ation, and plastic deformation. Theirmodel used thermal and
mechanical con-ditions to simulate the ultrasonic welding
process and led to the theory that ultra-sonic bonds are formed
due to high levelsof localized strain, high temperatures,
andplastic deformation along the interface.Siddiq et al. (Ref. 13)
also developed athree-dimensional model focusing on fric-tion and
heat generation at the interface.Their simulation determined the
effect offriction at the interface to be only useful inremoving
oxides and contaminates whilethe plastic deformation of material
actu-ally leads to a bond.
Multiple material combinations havebeen studied, including
aluminum, cop-per, titanium, and nickel (Refs. 11, 13, 14),as well
as many different fibers have beensuccessfully embedded including
fiber op-tics, silicon carbides, shape memory alloys,and
thermocouples (Refs. 3, 6, 9, 11, 15,16). The UAM process has been
found toeasily accommodate these embeddedfibers as the ultrasonic
energy allows forexcellent matrix material flow around
thefiber.
In all the above work, a one-to-one cor-relation of tensile and
shear propertieswith the underlying microstructure hasnot been
documented. Therefore, in thiswork, the mechanical properties of
alu-minum builds were measured and ob-served properties were
correlated with thedetailed microstructure evaluation usingoptical
microscopy, hardness mapping,and electron microscopy. The results
willbe compared with published literature onUAM processes as well
as data from ultra-sonic spot welding. The methodology anddata
generated in this research are ex-pected to provide a baseline for
the devel-opment of a very high power UAM (VHP-UAM) instrument
(Ref. 17). Thisinstrument will be capable of joininghigher-strength
alloys including titanium,copper, nickel, shape memory alloys,
car-bon steels, and low-alloy steels.
Experimentation
Alloys
In this research, a non age-hardenableAl-3003
(Al-1Mn-0.7Fe-0.12Cu wt-%)
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A B
Fig. 1 Schematic illustration of the UAM process. A Adding a new
tape layer; B periodic milling operation to form final
dimensions.
Table 1 UAM Processing Parameters Used in the Current
Research
ID Force Speed Amplitude Frequency Build Temperature(N) (mm/s)
(m) (kHz) (C)
Tack 350 59.3 12 20 150Weld 1150 42.3 17 20 150
Fig. 2 The UAM tensile and shear sample dimen-sions. A Reduced
shear specimens (1 and 2); B symmetric shear specimens (3 and 4); C
transversetensile specimens; D longitudinal tensile specimens.
A
C D
B
98.3
3.18
31.75
R 3.18
3.28
3.81
12.7 9.52
Tape LayerOrientation
Key
3.56
R 239
9.52R 3.18
9.52
17.14
6.3512.7
38.19.52
9.5238.1
17.14
12.7
38.7
38.1 9.52
9.52
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alloy was used as both tapes (H18, 150 mthick, 25.4 mm wide) and
substrate (H14,more than 12.7 mm thick). The composi-tion of the
materials used meets the stan-dard specification of the alloy (Ref.
18).
UAM Process Parameters
The ultrasonic sonotrode was madefrom Ti-6Al-4V alloy, and the
surface wassubjected to electrical discharge machin-ing (EDM) to
achieve the desired surfacetexture (Ra = 7 m). This surface
textureis known to provide consistent bond qual-ity (Refs. 5, 7).
During the tacking andwelding passes, the substrate was pre-heated
with a hot plate to 149C (300F)and was maintained at that
temperature.The preheat was used to soften the mate-rial, which
leads to better bonding. How-ever, during processing, the tape and
in-terface temperatures are not necessarilymaintained at this
preheat temperaturedue to complex heat transfer across themany weld
interfaces, heat generated atthe interfaces, and a heat-sinking
effectdue to the sonotrode.
Sequential joining of tapes to build asmall block was achieved
through tackingand welding passes. The differences be-tween the
tacking and welding passes are re-lated to the magnitude of the
process pa-rameters, i.e., normal load, travel speed,and amplitude
of ultrasonic vibration. Inthe current research, the vibration
fre-quency was kept constant at 20 kHz for allpasses due to machine
and sonotrode de-sign. Table 1 provides an overview of
theprocessing parameters used in the currentresearch. These
processing parameterswere obtained by extensive trial and
errorexperiments. One method of testing processparameters involves
joining of tapes by dif-ferent process parameters and manuallypeel
testing the builds. The best processingparameters are qualitatively
selected whenthe manual peel test fractures the tape,rather than
peeling off from the interfacial
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Table 2 Shear Test Results, Base Metal: Al 3003-H18 USS Is 110
MPa
Sample Force (N) Area (mm2) USS (MPa) % of BM
1 5089 81.6 62.4 56.72 4395 80.5 54.6 49.63 8830 215 41.1 37.44
11387 216 52.6 47.8
Average 52.7 47.9Standard Deviation 8.78 7.9
Stdev/Avg 0.167
Fig. 3 Sample image to demonstrate methodology used for image
analyses to derive the linear weld density. A Original optical
microscopy image; B processed image using ImageJ software.
BA
Fig. 4 Schematic illustration of steps to prepare TEM samples
from builds made with UAM. Sam-ples were taken along interfaces at
various heights (top, middle, and bottom) of the build. First,
opti-cal microscopy samples were prepared to select the regions of
interest. In the next step, the sample wastransferred to FIB
instrument. Then, an interface of interest was selected, and a
rectangular region oneither side of the interface was coated with
platinum. After this step, the focused ion beam machin-ing was made
on either side of the coated region. This leads to a thin film
sample that contains thebonded interface. In this schematic
representation, the n and n + 1 correspond to the successive
tapesduring the UAM processing.
Table 3 Transverse Tensile Tests, Base Metal: Al 3003-H18 UTS Is
200 MPa
Sample Force (N) Area (mm2) UTTS (MPa) % of BM
1 907 31.9 28.4 14.22 979 31.4 31.1 15.63 930 32.4 28.7 14.44
1010 31.4 32.1 16.15 859 33.5 25.7 12.96 601 31.4 19.1 9.67 1080
32.4 33.3 16.7
Average 28.3 14.2Standard Deviation 4.81 2.4
Stdev/Avg 0.170
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area. It is important to note that the processparameters used
here may not be optimumand are considered as the starting point
forthis and future research. Details of the peeltest instrument and
technique have beencovered extensively by other researchers(Refs.
1921).
Mechanical Property Testing
Previous mechanical strength studies on
UAM samples focused on peel tests (Refs.1, 2, 4, 6, 2224). Peel
tests, while useful forcomparison between parameter sets andother
UAM samples, are primarily used formeasuring adhesive strength of
tape, glue,or other bonded surfaces and do not pro-vide strength
values useful for the design ofbulk UAM parts. In order to be
utilized asan additive manufacturing process, bulkmechanical
strengths such as ultimateshear and tensile strengths must be
known
for design of UAM samples. To date, therehas been no reported
research on such bulkstrength properties. In order to obtain
bulkstrength properties of the UAM matrix,three types of samples
were made: lapshear, transverse tensile, and longitudinaltensile.
The geometries of these test speci-mens are presented Fig. 2. The
shearspecimens were built such that the tape in-terfaces were along
the shear plane. Sheartests were conducted using a specializedshear
jig and a compressive load with an av-erage displacement rate of
0.28 mm/s.
Initial shear test specimens had a re-duced interface area to
ensure failurebelow the 5000-lb machine capability. Ini-tial
estimates for strength assumed theshear strength for UAM specimens
wouldbe approximately 75% of the bulk mate-rial. As testing
revealed, the shearstrength was much lower than anticipated,and
later samples were not prepared witha reduced interface area. The
transversetensile specimens were built such that the
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Table 4 Longitudinal Tensile Test Results, Base Metal: Al
3003-H18 UTS Is 200 MPa
Sample Force (N) Area (mm2) ULTS (MPa) % of BM
1 2630 11.71 225 112.52 2900 12.21 238 119.03 2880 12.03 240
120.04 2870 12.13 237 118.55 2790 11.95 233 116.5
Average 234 117.0Standard Deviation 5.89 2.9
Stdev/Avg 0.025
A B
Fig. 5 Fracture surface image analyses of transverse tensile
test samples. A Low-magnification optical image; B
high-magnification optical image show-ing regions I and III after
processing. Region I is well bonded material with recrystallization
across the interface. Region II is deep valleys carved by the
sonotrodeduring the previous pass (not shown). Region III material
is directly opposite Region II and is unaffected during the UAM
process; C SEM image of the topsurface of the tape; D SEM image of
the bottom surface of the tape; and the featureless gray regions
are Region III material.
500 m
Top Surface Bottom Surface
500 mC D
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tape interfaces were perpendicular to theapplied axial force.
Transverse tensile testswere conducted using specialized speci-men
shoulder grips with an average dis-placement rate of 0.32 mm/s.
Longitudinaltensile specimens were built with tape in-terfaces
parallel to the applied axial forceand were tested using pinned
grips with anaverage displacement rate of 0.52 mm/s.
For all tests, samples were placed in auniversal
tension/compression testingframe and were stressed until failure.
Theapplied force was recorded using a ten-sion/compression load
cell and frame ac-tuator displacement with an integratedlinear
variable differential transformer(LVDT). Maximum loads were used to
ob-tain ultimate stresses, and the shape of theforce-displacement
plots was used to helpcharacterize specimen failures. Becausethe
integrated LVDT measures the testingframe actuator displacement,
all displace-ment data includes displacement gener-ated within the
load train as well as thespecimen. For this reason, the shape of
theforce-displacement plots can only be usedto determine if a given
sample failed in abrittle or ductile mode through
qualitativeanalysis. However, this cannot be used tocalculate
specimen strain or related prop-erties such as the elastic modulus.
Afterthe mechanical testing, the fracture sur-faces of the shear
and transverse tensilesamples were examined with optical
andscanning electron microscopy (SEM).
Optical Microscopy and HardnessMapping
Optical metallographic samples wereprepared using standard
metallographictechniques. The samples were preparedfrom cross
sections perpendicular to thetravel direction. Five optical images
at10 magnification were taken from dif-ferent locations within the
build. Eachimage corresponded to 1111 by 833 m,containing five
interfaces. These mi-croscopy images were analyzed with thepublic
domain ImageJ software program
(Ref. 25). With linearintercept analyses, theLWD was measured
asa function of distancein a direction perpen-dicular to the
metallictape layers. Grayscaleimage threshold val-ues (0 to 60)
were keptconstant to delineatethe void areas in allthese images. A
typicaloptical image beforeand after thresholdprocessing
demon-strates the effective-ness of delineating thevoids between
layers Fig. 3A, B.
For the microhard-ness testing, a LecoAMH-43 machine wasused to
create a 200 20 map of hardnessindents with a diamond indenter.
Themeasurements were made with 25-g loadand a 13-s dwell time, and
spacing be-tween the hardness indents was 150 m inboth directions.
The coordinates of theindents were designed to sample the
solidmatrix regions away from interfaces.Hardness measurements were
made ontapes that were not ultrasonically consol-idated in the same
orientation, as areference point.
Analytical ElectronMicroscopy
In order to examine the grain structureand morphology in
specific locations (bot-tom, middle, and top regions of the
build)through transmission electron microscopy(TEM), the samples
were prepared usinga FEI Helios dual-beam focused ion beam(FIB)
microscope. The samples were pre-pared from cross sections
perpendicularto the travel direction along interfaceswith apparent
good bonding. The FIB
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Fig. 6 Measured load vs. displacement curves. A Shear tests with
sym-metric and reduced cross sections; B transverse tensile tests;
C longitu-dinal tensile tests.
A B
C
Table 5 Linear Weld Density (%) from Optical Micrographs Taken
from Random Locationswithin the Build
Image Interface StandardNumber 1 2 3 4 5 Average Deviation
1 49.5 86.2 74.7 91.1 48.2 69.9 20.22 70.4 74.1 61.9 88.1 51
69.1 13.83 69.5 59.6 67.5 65.5 5.24 53.1 65.7 52.2 55.2 46.7 54.6
7.05 34.2 65.6 68.1 91.6 75.2 66.9 20.9
Overall 65.2 15.3
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contains both an electron beam as well asan ion beam that can be
used for imaging.The electron beam is a standard second-ary
electron beam, which can be used toimage topological difference but
does notreveal grain structure in unetched alu-minum alloys. The
ion beam was used toimage the grain structure of the materialwith
contrast differences arising from gal-lium ion channeling contrast.
To createTEM foils, platinum is deposited over theregion of
interest to protect the foil sur-face from Ga+ implantation
duringmilling. Trenches are then milled on bothsides of the
platinum to create the foil.The sample is then bonded to an
om-niprobe needle also using platinum. Thesides of the foil are
then milled to create afree-standing foil. Once the sample is
cutfree, it is lifted out using the omniprobeneedle and welded to a
copper grid usingplatinum. Once the sample is welded tothe grid, it
is thinned using ion milling anda series of various apertures at 30
kV. Thesteps used in making the samples areschematically
illustrated in Fig. 4. Finally,the samples were then examined using
aFEI Tecnai F20 operated in STEM mode.
Results and Discussions
Mechanical Properties
The original mechanical properties ofAl 3003-H18 alloys are as
follows: The ul-timate tensile strength (UTS) is 200 MPa,the yield
strength (YS) is 186 MPa, andthe ultimate shear strength (USS) is
110MPa (Ref. 18). The mechanical propertydata from this research
program are sum-marized in Tables 24. All shear tests re-sulted in
a linear force-displacement re-lationship, indicating samples
failed in amacro-level brittle fracture mode. Asshown in Table 2,
an average USS of 52.7MPa with a standard deviation of 8.78MPa was
found. The average USS was ap-proximately 48% of that of the solid
basematerial. The results from the transverse
tensile tests are shown inTable 3. The average ulti-mate
transverse tensilestrength (UTTS) was 28MPa, approximately 15%of
the tensile strength ofsolid base material. Stan-dard deviation for
UTTSwas 4.57 MPa.
To understand the re-duced strength of thetransverse tensile
testsamples, the fracture sur-faces were characterizedusing optical
and scan-ning electron microscopy.The fracture surfacesfrom the
samples from 1to 5, as well as samples 7and 8, share a similar
frac-ture surface Fig. 5.These images indicatethat the interface
regionshave many small speckled-like features dispersedthroughout
the bond area. This feature isdue to small areas of bonded
material(marked as I in Fig. 5A) mixed with smallareas of
unaffected material (marked asIII in Fig. 5B). In order to make
sure theRegion I (Fig. 5B) is a true bond, the frac-ture surfaces
from either sides of the frac-ture were characterized with
scanningelectron microscopy (Fig. 5C, D). Scan-ning electron
microscopy showed that thefracture surfaces do show localized
duc-tile failure with typical microvoid coales-cence features. This
is irrespective of thefact that the load-displacement curves donot
show appreciable macro level ductil-ity. These observations proved
that theultrasonic additive manufacturing didnot reduce the
inherent ductility of thematerial; however, on a macro scale,
thematerial behaved in a brittle fashion dueto the voids. The lack
of a yield point andhardening region is related to prematurebrittle
fracture caused by the voids.
Transverse tensile samples behaved sim-ilar to the shear
samples, as indicated by the
linear force-displacement plots Fig. 6A,B. From Fig. 6B, sample
6 is considered tobe an outlier, as seen by its much lower fail-ure
force when compared to the other sam-ples. Upon examination, the
fracture sur-face of the sample 6 showed interestingfeatures that
were different from the othersamples inspected Fig. 7A. Optical
mi-croscopy showed trenches and ridges, whichare typical of
surfaces created by a millingoperation (Ref. 7). Furthermore,
scanningelectron microscopy showed that the bondshave formed along
these ridges and havefailed again by ductile mode Fig. 7B,
C.Cursory evaluation of the above fracturemorphology may be
puzzling; however, thisphenomenon may be explained. During theUAM
process, at frequent intervals amilling operation is performed to
achieve aflat surface to ensure dimensional accuracyof the finished
part. In sample 6, the failureoccurred at one such milled
interface. Thefracture surface showed that the area frac-tion of
the bonded region was small com-pared to unbonded regions, likely
the rea-
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A
Fig. 7 Fracture surface analyses of transverse tensile test 6. A
Opti-cal image showing crosshatching from the milling pass; B SEM
imageof the top surface of the tape containing fractured regions.
It is apparentall bonding occurred in the light-colored regions,
less than half of theavailable surface area; C SEM image of the
bottom surface of the tapecontaining fractured regions. It is
apparent all bonding occurred in thelight-colored regions, less
than half of the available surface area.
B
C
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son for the premature failure of this sample.All of the samples
tested had this flat passwithin them; however, only sample 6
failedin this manner. Because it is not known whythe other samples
did not fail at this loca-tion, it is believed that sample 6 does
notgive a true representation of UAM bondtensile strength and is
also excluded in sta-tistical analysis. The sensitivity of
surfaceroughness on the bond quality has been ad-dressed by
previous researchers (Refs. 7,22) by relating the surface roughness
of thesonotrode to changes in linear weld density.
Unlike the other UAM samples, thelongitudinal tensile samples
(Table 4) ex-hibited a substantial plastic yielding re-gion after
the linear elastic region Fig.6C. This is more typical of aluminum
al-loys and indicates that failure occurred ina ductile mode. All
tested samples exhib-ited a higher than expected tensilestrength.
The average ultimate longitudi-nal tensile strength (ULTS) was 234
MPa,17% more than the original Al 3003-H18tape based on published
properties (Ref.18). This is a departure from both thetransverse
tensile samples and shear sam-ples previously tested in which the
failurestresses were significantly lower than theAl 3003-H18 tape
based on publishedproperties(Ref. 18). In this orientation, nodrop
in tensile strength was expected asthe load was transmitted along
the solidtapes as opposed to across the interfacesbetween them.
However, the increase instrength above the base material was
notexpected and further explanation of thisphenomenon is
required.
Microstructure and MechanicalHeterogeneity
To rationalize the reduction in me-chanical properties in the
transverse load-ing condition, the LWD of the builds indifferent
regions were analyzed. A typicaldata set of linear void density
(inverse ofLWD) is shown in Fig. 8. The image analy-ses show the
LWD can vary from 35 to99%, depending on the interface. The
av-erage LWD of all the images analyzed wasfound to be 65.215.3%,
(Table 5). Imageanalyses of optical micrographs of thefracture
surface of the transverse tensilesamples (Fig. 5) yielded 662%
bondedarea. In Fig. 5D, it is clear the voided re-gions are random
in nature. In stereologi-cal terms, randomly placed line segmentsin
cross-sectional images are proportionalto an objects area in a 2-D
plane (Ref. 26).However, in the current study, only
onecross-sectional plane was used. This pre-vents the conclusion
that LWD is directlyrelated to area density of properly
bondedmaterial in UAM builds, despite the aver-ages being
comparable. With additionalangular cross sections and more
samples,it may be possible to confirm a possible
one-to-one relationship. Regions withlower amounts of bonded
area within thebuild are expected to reduce the trans-verse and
shear strength significantly. Thishypothesis is consistent with the
conclu-sions made by previous researchers thatlinear weld density
is a good measure ofUAM bond quality.
The mechanical properties measuredalong longitudinal sections
showed a 17%
increase in ULTS compared to that of theoriginal Al 3003-H18
tape materials. Inorder to rationalize this increase instrength,
hardness mapping was per-formed on the UAM builds. The mapshows a
soft substrate and UAM build re-gions with large variations in
hardness Fig. 9A. The hardness data were analyzedin terms of
frequency distribution Fig.9B. This graph also shows the
hardness
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Fig. 8 A Plot of linear void density (LVD) vs. five interfaces;
B the corresponding optical imagenumber 5. The LVD point for each
interface was taken as the high point and is shown. Linear void
den-sity is the inverse of linear weld density.
Fig. 9 Microhardness plot of a UAM build. A The map was 200
indents tall by 20 indents wide, with thesofter substrate at the
bottom. No gradient in hardness (either from bottom to top of build
or left to right ofbuild) was observed. This indicates later passes
have minimal effects on the hardness of previously depositedlayers;
B histogram showing bimodal hardness distribution with the UAM
build foils significantly harderthan the substrate. The hardness of
unconsolidated foils is also overlaid on same plot and is below the
peakhardness of the UAM build; C optical image showing a high
hardness foil indent on left (90 HV) next to aweaker interface
indent on the right (30 HV). Interface areas with voids or defects
caused by insufficient ma-terial flow to fill in grooves cut by
sonotrode during previous pass had lower hardness.
A B
B
Linear Void Density (%)0 10 20 30 40 50 60 70 80
0
200
400
600
800
1000
1200
1400
24.7
8.5
31.9
33.8
65.8200 m
Imag
e H
eigh
(
m)
C
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distribution from the original Al 3003-H18 foils, which had an
average hardnessof 64.52.7 HV. The UAM build had anaverage hardness
of 73.71.9 HV. Thedata show that the UAM builds are indeedharder
than the stock foils and provide aqualitative explanation of the
increase inULTS. Careful analyses of hardness in-dents in certain
regions also showed inter-esting features Fig. 9C. In one region,
asmall indent showing high hardness wasright next to a large indent
showing lowhardness. This low hardness was associ-ated with a large
planar defect (marked byarrows). Although the weakened regionsmay
be explained with the presence of un-bonded areas, it is necessary
to evaluatethe hardened regions through detailed mi-crostructure
characterization.
Transmission electron microscopesamples from the interfaces from
the bot-
tom (near substrate), middle, and top re-gions of the build were
extracted throughFIB machining. The electron microscopyimages are
presented Fig. 10AE. Themicrostructure of the original foil is
alsoprovided for comparison Fig. 10F. In-terestingly, the
microscopy images fromthe bottom (Fig. 10A) and middle (Fig.10B)
regions failed to show any sharp in-terface region indicating the
formation ofa metallurgical solid-state bonding. Thegrains were
equiaxed in nature, quite dif-ferent from that of elongated grains
of theoriginal Al 3003-H18 tapes Fig. 10F.This suggests that the
bond formation maybe associated with recrystallization. In
ad-dition to the equiaxed grains, fine Al-Mn-Fe-based
intermetallics were observed inthe samples along the grain
boundariesand within the matrix grains. These inter-metallics are
found in the original Al 3003-
H18 tapes and do not appear to be af-fected by the UAM process.
The interfacemicrostructure (Fig. 10C) from the top re-gion showed
interesting features. Theoriginal interface location can be
inferredfrom the sudden change in the grain struc-ture. The
microstructure in the (n + 1)thtape shows the original pancake
structure,which transitions sharply to a coarse andrecrystallized
grain structure close to theoriginal interface location. The
mi-crostructure from the nth layer does notshow any pancake
structure, rather morerecrystallized structure. Moreover, a re-gion
of grain boundary decohesion wasalso observed. This grain boundary
deco-hesion was also confirmed with high-mag-nification analyses
Fig. 10D. A surveyof many samples from different regionsalso showed
the interface regions con-tained fine recrystallized grains (<
500
MAY 2010, VOL. 89112-s
WELDING RESEARCH
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C
D
B
Fig. 10 Six TEM images. A Bright field TEM image taken from an
interface location with apparent good bonding. The interface cannot
be determined easily,indicating potential recrystallization across
the interface. Small, white Al-Mn-Fe intermetallics can be seen
here; B another interface location again showing the dif-ficulty in
discerning the bond line; C a third interface location where the
bond line can be determined, as pointed out by the red arrows. The
blue arrow points to asmall void that appears to have migrated from
the interface into the bulk of the material; D high-magnification
bright field image of the void in C; E dark fieldimage showing the
high levels of dislocations within the grains and the Al-Mn-Fe
intermetallic particles; F bright field image of the original foil
before consolida-tion. The as-rolled structure is pancake-like
grains with some dislocations present. Dislocation content
significantly lower than that observed after UAM processing.
E F
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nm) with relatively low dislocation densityand coarse grains
(500 nm to 2 m) withrelatively high dislocation density asshown by
dark field microscopy Fig.10E. The original foil, before
consolida-tion, shows (Fig. 10F) pancake-like grains,as expected
from as-rolled material. Dis-locations were present, though a
muchlower concentration than in the grainsalong interfaces after
UAM processing.
Discussion on Process-Structure-Property Correlations
In order to understand the interfacemicrostructure, it is
important to reviewthe steps involved in the UAM process,shown
schematically in Fig. 11. In Fig.11A, a first layer has been bonded
to thesubstrate, with the top of this layer left ina rough
condition after the sonotroderolled over it. When the next layer is
ap-plied, the bottom of the new layer is rela-tively flat, creating
an interface between asmooth surface and a rough surface Fig.11B.
When the sonotrode comes directlyon top of the interface during the
tackingpass, the relative motion between the twolayers creates
frictional and deformationalheating and partially collapses
asperities
Fig. 11C. This results in a weak bondbetween the layers, with
many voids asshown in Fig. 11D. During the weldingpass, more
ultrasonic energy (higherforces and amplitude of vibration) is
usedto finish the bond Fig. 11E. Some resid-ual voids remain, as
shown in Fig. 11F.
The final microstructure at the inter-face can be summarized to
consist of threeregions as shown in Fig. 11G. During thewelding of
the previous layer, the top sur-face of each foil interacts with
thesonotrode and becomes rough. This roughsurface becomes the
bottom of the follow-ing interface. Where peaks occurred alongthe
rough surface, contact was made withthe next foil and a bond
resulted (RegionI). This region constitutes
recrystallizedmicrostructure (500 nm to 2 m) acrossthe interface
and has good metallurgicalbonding. It is believed when these
peaksare brought into contact with the new foillayer sufficient
strain energy, tempera-ture, and forces exist to force dynamic
re-crystallization. However, where valleys oc-curred due to the
sonotrode texture,Region II, they were often too deep tomake
contact with the next layer beingadded. This resulted in voids
along the in-terfaces and created the Region III mate-
rial on the foil directly above it. Region IIIis the unaffected
original foil surface thathas not been touched by either
thesonotrode or the foil layer beneath it. Re-gion III was only
found on the top surfaceof the interface. Region II material was
di-rectly opposite and the cause of Region IIImaterial. In this
study, focus was given tounderstand the mechanism of the
grainstructure evolution in Region I. Based onthe microstructure
from Fig. 10AC andE, we can conclude that the original pan-cake
grain structure was modified to formsub grains with sizes ranging
from 500 nmto 2 m with different levels of dislocationdensity. To
understand this reduction ingrain size, we assume this process is
simi-lar to that of hot working of aluminum al-loys. The subgrain
size (dsub in m) duringhot working can be related to
Zener-Hol-lomon (Zh) parameter and peak tempera-ture (TP) achieved
during hot working(Refs. 27, 28).
Equation 1 has been used to estimatethe grain size in both
friction and frictionstir welding. The Zener and Hollomon(Zh)
parameter has been estimated for
d log Zsub h = +[ ]0 60 0 08 11. . ( ) ( )
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H
A
D E F
G
B C
Fig. 11 Schematic representation of the UAM process highlighting
the vari-ous stages. A Beginning of new layer, top of previous
layer textured bysonotrode during previous layer bonding. B New
layer, 2, placed by feedingmechanism in front of sonotrode; C
sonotrode tacks the new layer down, gen-erating frictional heat and
forming a weak bond; D new layer tacked down,many residual voids
present; E sonotrode passes again for welding pass, de-forming the
top surface as it passes; F layer 2 attached, some voids are
stillpresent between layers 1 and 2; G third layer ready to be
added. Enlargementof bond interface showing the three regions.
Region I is well-bonded material,Region II is valleys carved by
sonotrode, and Region III is untouched material.
Fig. 12 A contour plot estimating the final grain size using the
Zener-Hol-loman parameter. The numbered, curved lines represent the
possible combi-nations of temperature and strain rate to achieve
the same grain size. Dashedlines represent the range of grain sizes
found along UAM interfaces. The ver-tical line is the calculated
strain rate derived from Equations 36.
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aluminum alloys as a function of strainrate ( in s1) and peak
temperature (TP)(Ref. 29).
In order to understand the subgrainstructure in UAM process,
Equation 2 wasused. The strain rate during ultrasonic ad-ditive
manufacturing is calculated usingthe following approximation. The
totaldisplacement due to the plastic deforma-tion, a thin slab of
material under thehorn, can be taken as the horn amplitude,i.e., (
d = 26 106 m). This assumesthere is no slippage of the interface
mate-rial. The asperity height is estimated as thepeak-to-peak
height of the tape surface.This surface is assumed to be a
negativeimage of the sonotrode texture, which hasa value of 7 106
m, as reported by John-son (Ref. 7). Thus, the peak-to-peakheight
of the average asperity is 14 106m. Furthermore, the height of the
asperi-ties is assumed to have negligible changewith respect to
time. With these assump-tions, displacement of the bonded
regionswith respect to time can be given by theexpression:
Asperity velocity is calculated as the de-rivative of
displacement with respect totime:
The shear strain of an asperity is given bythe equation:
Shear strain rate is then found by takingthe derivative of
strain with respect to timeas follows:
Over one ultrasonic cycle, an asperitywill have a strain rate
varying between2.3 105 rad/s with an RMS value of
1.1 105 rad/s. Because we do not knowthe peak temperatures
experienced by in-terface regions, the subgrain sizes wereevaluated
as a function of peak tempera-ture and strain rate. This is shown
as aform of contour plot Fig. 12. The cal-culated micro strain rate
from Equation 6results in a peak temperature of around300 K for the
500-nm grains and a peaktemperature of 900 K for the 2-m
grains.This range of temperatures is larger thanexpected, but this
may have been causedby the approximations in Equations 1 to 6.In
Equation 6, perfect transfer of strainwas assumed, no slipping
between thesonotrode and the foil was accounted for.Account for
slipping, the resulting strain,and therefore peak temperature
requiredto achieve a certain grain size, would havedecreased.
Meanwhile, Equations 1 and 2were developed for simple monotonoushot
working conditions and not reversiblestrains that are experienced
during UAM.Gunduz et al. (Ref. 30) estimated a localstrain rate of
1 104 s1 at a temperatureof 513 K based on vacancy calculationsand
the diffusion profile observed in alu-minum-zinc ultrasonic welds.
This workwas based on finding the vacancy concen-tration required
to reduce the meltingtemperature so the observed small meltregion
was possible at ultrasonic weldingtemperatures. Their result is
within therange of grain size, temperatures, andstrain rates
studied here. Conversely,macro strain rates were studied by Gaoand
Doumanidis (Ref. 31) by placing astrain gauge near, but not
directly be-neath, the welding sonotrode. They foundmaximum strains
of 90 106 over 0.5 s or1.8 104 s1. This low strain rate is
ex-pected as Gao and Doumanidis measuredmacro strains with a strain
gauge of amuch larger size scale than the asperitiesused in
Equation 6.
Recently, Johnson has proposed thatthe materials under
reversible strainingconditions may exhibit an UltrasonicBauschinger
effect (Ref. 7). However,the interaction of these effects with
heat-ing and subgrain formation is not clear.In addition, the
estimated strain rateshave to be validated based on detailed
fi-nite element deformation models (Ref.12), which considers the
spatial variationsas well as dynamic strain hardening orsoftening.
The localized temperaturealong interfaces may be affected by
thefriction and rapid deformation condi-tions. In the current UAM
process, thesubstrate temperature is maintained at149C (422 K).
This heat will diffuse fromthe substrate to the entire build. As a
re-sult, with the progress of UAM builds, thepreviously welded
interfaces will be sub-jected to an isothermal hold close to
thistemperature throughout the processingof the build. This
isothermal hold is also
expected to induce some of the recrystal-lization and grain
growth observed. Thissuggests the need for measuring the spa-tial
and temporal variations of the tem-perature during the UAM process.
Thiswill be the focus of the future work (Ref.32). The next step is
to provide some di-rections to rationalize the measured me-chanical
properties. From the above dis-cussions, it is apparent that all
UAMsamples will have large voids along the in-terfaces as well as
localized hard and softregions. The voids can be treated as
em-bedded cracks, which cause stress con-centrations that resulted
in brittle frac-ture of the shear and transverse tensilespecimens.
The loading of the transversetensile samples results in a mode I
frac-ture, while the shear sample loading in-duces mode II
fracture. For a given crackand load magnitude, mode I
fractureloading typically exhibits the largeststress intensity
factor (SIF) followed bymode II fracture (Ref. 33). Because
load-ing is parallel to the embedded cracks inthe longitudinal
tensile samples, there isno SIF and the strength in this
orienta-tion was not reduced. This fracture me-chanics perspective
further explains whythe shear and transverse tensile sampleshave
lower than normal strengths andbrittle fracture characteristics,
while thelongitudinal samples were not weakenedby the presence of
voids and insteadfailed with ductile characteristics. Again,the
above discussion is simplistic, doesnot provide a predictive
capability, anddoes not account for all transients thathave been
observed, such as the varia-tions in tensile testing shown Fig.
6.Further work is necessary to develop de-tailed computational
models that incor-porate the spatial variation of mi-crostructure
and voids and constitutiveresponse of the bulk and interface
loca-tion. To facilitate the development ofconstitutive properties
of the interfacelocations, the grain size distributionalong the
interface has to be character-ized close to the voids and away from
thevoids using orientation-imaging mi-croscopy. This grain
orientation and sizedistribution will allow us to develop
mul-tiscale models similar to the ones beingdeveloped by Ghosh
(Ref. 34).
Finally, in order to overcome the deepchannels carved by the
sonotrode, a veryhigh power UAM system is being developedby EWI
(Ref. 17). It is believed that higherultrasonic power input, higher
amplitudes,and normal forces will increase the plasticflow at the
interfaces. This should enablegreater LWD, reducing the inherent
stressconcentrations and improving the tensileand shear strength of
UAM builds. Higherplastic flows should also improve the
metal-lurgical bonds by ensuring all of the oxidesare removed from
the interface.
Z expTh p
=
18 7722
,( )
( )( )
td th
d
1
=
+
1
2(( )
( )
t
h
d th
1
=
+
1
2
v t
h
( )( )6
=
tand t
h1 5
( )( )
d t v t cos t
= = ( ) ( ) . ( 20,000 )
3 3 2
(4)
d t sin t( ) ( 20,000 )
= 26 10 26 (3)
MAY 2010, VOL. 89114-s
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Conclusions
The present study focused on linkingmicrostructure and LWD to
mechanicalproperties of ultrasonic additive manufac-turing builds.
Using TEM, SEM, and op-tical microscopy along with microhardnessand
tensile and shear testing, the micro-scopic and macroscopic
properties ofUAM builds were analyzed. The followingwas found:
1. The average shear strength of thetested UAM samples was
approximately48% of the expected 110 MPa ultimateshear strength of
Al 3003-H18. The aver-age transverse tensile strength was
ap-proximately 14% of the expected 200 MPatensile strength of Al
3003-H18. Trans-verse tensile and shear testing results
areindicative of bond quality alone; failureoccurs before
microstructure becomes significant.
2. Without optimized parameters,UAM weldments result in voids
scatteredthroughout all interfaces. This ultimatelycaused the
samples to fail in a low ductil-ity manner with low strength
values.
3. Image analysis of cross-sectionedsamples found an average
linear weld den-sity of 67.416.1%. Image analysis oftransverse
tensile fracture surfaces foundan average area weld density of
662%. Adirect comparison between LWD and areaweld density was not
possible based on thesample size.
4. The average longitudinal tensilestrength was approximately
117% of theexpected tensile strength of Al 3003-H18.This indicates
the foils were strengthenedduring processing and was confirmed
bymicrohardness testing. Microhardnesstesting found the average
hardness of theUAM foils increased almost 15%, from64.52.7 HV in
the original foils to73.71.9 HV, during processing.
5. A hypothesis relating grain refine-ment to strain and
temperature using theZener-Hollomon parameter was devel-oped.
Microstrain rates were estimatedbased on operating conditions to
bearound 1 105 s1. From this and an ob-served grain size of 500 nm
to 2 m, an es-timated peak temperature range for theUAM process of
300 to 900 K was calculated.
Acknowledgments
The authors would like to thank the Co-operative Research
Program of EdisonWelding Institute for supporting this re-search.
In addition, we thank Dr. K. Graff(EWI), Dr. M. Sriram (OSU), and
MattShort (EWI) for suggestions and fruitfuldiscussions during
preparation of themanuscript. R. M. Hahnlen and M. J.
Dapino are grateful to the member organi-zations of the Smart
Vehicle Concepts Cen-ter (www.SmartVehicleCenter.org), the
Na-tional Science FoundationIndustry/University Cooperative
ResearchProgram (I/UCRC), and the Smart VehicleConcepts Graduate
Fellowship Program.
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