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\begin{document}

\title{Microstructural and fatigue crack growth behavior inhomogeneity of trace
Sc added 7085 alloy friction stir welded joint}



\author[1]{Zixing Jia}%
\author[1]{Lili Wei}%
\author[1]{Hongfeng Huang}%
\author[1]{Chongyu Liu}%
\author[1]{Hongjie Jiang}%
\author[2]{Xuda Xu}%
\affil[1]{Guilin University of Technology}%
\affil[2]{Sanshui Fenglv Aluminium Co., Ltd.}%


\vspace{-1em}



  
  \date{May 27, 2024}


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\selectlanguage{english}
\begin{abstract}
The inhomogeneity of friction stir weld (FSW) joint of a 7085 Al-based
alloy containing trace Sc was investigated by comparing the hardening
phase morphology, grain structure, mechanical properties and fatigue
crack growth (FCG) behavior in based metal (BM), heat-affected zone
(HAZ) and weld nugget zone (WNZ). The results showed that
microstructural evolution during FSW was remarkably affected by Sc
addition. The recrystallization and growth of grains in BM and HAZ were
restrained by coherent Al3(Sc,Zr) particles. In HAZ, partially
dissolution and slight growth of precipitates resulted in softening but
contributed to enhance the FCG resistance according to shear mechanism.
In WNZ, high frictional heating and intense plastic deformation induced
recrystallization and supersaturated solid solution which decomposed to
form new reprecipitates, including many new GPII zones, ' phases and
some coarse equilibrium phases attached to grown incoherent Al3(Sc,Zr).
Owing to combined influences of fine random orientated grain structure
and reprecipitates, WNZ possessed the best fatigue endurance.%
\end{abstract}%



\sloppy


\textbf{Microstructural and fatigue crack growth behaviorinhomogeneity
of trace Sc added 7085 alloy friction stir welded joint}

Zixing Jia\textsuperscript{2}, Lili
Wei\textsuperscript{1,2,}11\textsuperscript{*}Corresponding author 1.
Tel.: +86 13517739930. E-mail address:\emph{hhfwll@163.com}. (Lili Wei),
Hongfeng Huang\textsuperscript{1,2,}22\textsuperscript{*}Corresponding
author 2. Tel.: +86 13517738716. E-mail
address:\emph{csuworld@163.com}.(Hongfeng Huang), Chongyu
Liu\textsuperscript{2}, Hongjie Jiang\textsuperscript{2}, Xuda
Xu\textsuperscript{3}

\textsuperscript{1} Key Laboratory of New Processing Technology for
Nonferrous Metal \& Materials, Guilin University of Technology, Ministry
of Education, Guilin 541004, China

\textsuperscript{2} College of Materials Science and Engineering, Guilin
University of Technology, Guilin 541004, China

\textsuperscript{3} Sanshui Fenglv Aluminium Co., Ltd., Foshan, 528000,
China

\textbf{Abstract}

The inhomogeneity of friction stir weld (FSW) joint of a 7085 Al-based
alloy containing trace Sc was investigated by comparing the hardening
phase morphology, grain structure, mechanical properties and fatigue
crack growth (FCG) behavior in based metal (BM), heat-affected zone
(HAZ) and weld nugget zone (WNZ). The results showed that
microstructural evolution during FSW was remarkably affected by Sc
addition. The recrystallization and growth of grains in BM and HAZ were
restrained by coherent Al\textsubscript{3}(Sc,Zr) particles. In HAZ,
partially dissolution and slight growth of precipitates resulted in
softening but contributed to enhance the FCG resistance according to
shear mechanism. In WNZ, high frictional heating and intense plastic
deformation induced recrystallization and supersaturated solid solution
which decomposed to form new reprecipitates, including many new GPII
zones, ?' phases and some coarse ? equilibrium phases attached to grown
incoherent Al\textsubscript{3}(Sc,Zr). Owing to combined influences of
fine random orientated grain structure and reprecipitates, WNZ possessed
the best fatigue endurance.

\textbf{Keywords:}Aluminum alloy; Friction stir welding; Fatigue crack
growth; Mechanical properties.

\textbf{1. Introduction}

Al-Zn-Mg-Cu alloys (7xxx series) have been widely used in fields of
aviation, aerospace, and civilian transport owing to their high strength
to weight ratio, good fracture toughness and corrosion resistance
{[}1-3{]}. For the preparation of complex structural parts in
applications, the fusion welding technologies, e.g. Tungsten Inert Gas
process, Gas Metal Arc Welding are the most used methods to join these
high strength aluminum alloys. However, the joints made by these welding
method showed high hot cracking susceptibility in 7xxx series alloys
with high Cu content {[}4{]}. In addition to hot cracks, welding defects
like porosity, lack of fusion and incomplete penetration also frequently
took place. In order to avoid these problems associated with the
liquid-solid coupling, a newly invented process, friction stir welding
(FSW), which is a solid-state welding technique has been wildly used for
defect-free and efficient welding of 7xxx series alloys {[}5, 6{]}.

Unlike other conventional fusion welding methods, FSW is a complicated
thermal-mechanical process. Because of the combined influence of
frictional heating and intense plastic deformation, FSWed joints have
generally inhomogeneous complex spatial arrangements of microstructure
and mechanical properties {[}7-9{]}. In heat-treatable 2xxx and 7xxx
alloys, recovery/recrystallization of substructure and
re-dissolution/coarsening of hardening precipitates usually occur in
weld joint due to great frictional heat generated between welding tool
shoulder and workpieces {[}10{]}, resulting pronounced degradation in
mechanical properties as the absence of strengthening of dislocations,
refinement structure and precipitates {[}11, 12{]}. In order to decrease
these adverse effects, one approach is to reduce frictional heat, such
as increasing tool traverse speed/rotation rate ratio, cooling FSW joint
by water (underwater FSW). However, these low heat input FSW strategies
would raise the possibility of non-welded groove on the root side of the
weld {[}13, 14{]}. Another alternative is to increase the
recrystallization resistance of workpiece. It has been widely proved
that trace Sc can improve recrystallization temperature of aluminum
alloys due to the formation of
Al\textsubscript{3}Sc/Al\textsubscript{3}(Sc,Zr) precipitates in matrix
{[}15, 16{]}. The strength benefits of minor Sc addition have been
demonstrated in 7055 and 7075 alloys {[}17, 18{]}.

7085 aluminum alloy is a newer 7xxx alloy designed for aerospace
applications. It has been widely used in many aircraft structural
components due to high fracture toughness and low quench sensitivity
{[}19-21{]}. Because the Cu content of the 7085 alloy exceeds 0.8 wt.\%,
there are challenges in welding by using fusion welding technologies. As
the outstanding advantages of FSW technology, the application of FSW in
7085 alloy aroused widespread concern in recent years. Xu et al. have
investigated the FSW behaviors of 7085 alloy and discussed the effect of
welding parameters on the microstructure inhomogeneity and mechanical
properties of FSWed joint {[}5, 22-24{]}. However, because of the
increasing use of 7085 alloy in fatigue critical applications, fatigue
crack growth (FCG) behavior of the FSWed joint should be highly
regarded, but the related appropriate experimental data is insufficient
so far. Moreover, the influence of trace Sc addition on FSW behaviors of
7085 alloy has also rarely been reported, the effect of Sc on the
microstructural evolution and mechanical properties, especially FCG in
FSWed joint of alloy 7085 is not clear yet.

Therefore, in present work, a 7085-0.25Sc (wt.\%)-T6 Al alloy plates
were welded by FSW. The microstructures (sub-structures, precipitates)
and mechanical properties (hardness, tensile property and resistance to
fatigue) of base metal (BM), heat-affected zone (HAZ) and weld nugget
zone (WNZ) have been thoroughly investigated. The effects of Sc addition
on the microstructures and properties, especially FCG inhomogeneity of
FSWed joint were discussed in detail and elucidated.

\textbf{2. Experimental procedures}

2.1. Materials and FSW parameters

The 7085 with trace Sc alloy (hereinafter referred to as 7085Sc alloy)
used in this study was supplied as 4.5 mm-thick rolled plates. The
chemical composition is listed in Table 1. A solid solution treatment
(480\selectlanguage{ngerman}°C/1h, quenching in water) and aging annealing treatment (120°C/24h,
cooling in air) were applied to obtain T6 plate samples for welding.
Before FSW process, the surfaces of plates were polished to remove the
oxide layer for the purpose of minimizing welding defects in welding
joint. The friction stir pin tool had a 15 mm concave shoulder, and a
taper threaded probe whose diameter was 3.5 mm and long was 4 mm. The
welding direction was parallel to the rolling direction of the plates.
The weld travelling speed was 100 mm/min and ration rate of pin tool was
900 RPM.

\textbf{Table 1.} Chemical compositions (in wt.\%) of the 7085Sc alloy\selectlanguage{english}
\begin{longtable}[]{@{}llllllll@{}}
\toprule
Zn & Mg & Cu & Sc & Zr & Si & Fe & Al\tabularnewline
\midrule
\endhead
7.91 & 1.39 & 1.79 & 0.26 & 0.12 & 0.14 & 0.06 & In
balance\tabularnewline
\bottomrule
\end{longtable}

2.2. Microstructural analysis

The morphologies of grains and precipitates in BM, HAZ and WNZ were
investigated. Optical microscopy (OM) observation was performed using
Leica DFC 295 in polar light mode. Electron backscatter diffraction
(EBSD) analyses were performed using ZEISS Sigma 300 field emission gun
scanning electron microscope with a scanning step size of 0.15 \selectlanguage{greek}μ\selectlanguage{english}m. The
EBSD data was analyzed by Channel 5 software to acquire microstructure
statistics information. Boundaries with misorientations less than 15\selectlanguage{ngerman}°
were defined as low-angle grain boundary (LAGB) and those with
misorientations greater than 15° were defined as high-angle grain
boundary (HAGB), respectively. The scanning electron microscopy (SEM)
observation of fracture surface were also operated by ZEISS Sigma 300 at
20 KV. Transmission electron microscopy (TEM) and high resolution
transmission electron microscope (HRTEM) observations were carried out
on FEI Tecnai G\textsuperscript{2} F20.

2.3. Mechanical properties tests

The microhardness arrangement of the entire weld cross section was
detected by a LECO-AMH43 Vickers hardness tester with a testing load of
300 gf and a dwell time of 13 s. The spacing between adjacent
indentations was 0.5 mm. Small dog-bone shape tensile testing specimens
were machined by wire cut electrical discharge machining and tested on
AGS-X-50kN testing machine at a strain rate of
1.3×10\textsuperscript{-3} s\textsuperscript{-1}. The FCG rate tests
were performed on INSTRON-8801 testing machine at room temperature by
using the compact tension (CT) specimens. All fatigue testing specimens
were polished and tested with a load ratio of R=0.1. Sine wave loading
with a frequency of 10 Hz was used. The crack length was measured
continuously by the compliance method with crack mouth clip gauge
(3541-010M-040M-ST). The crack length and cycle number were recorded at
every increment of 0.05 mm in crack length. Fig. 1 shows the dimensions
of CT specimens and the sampling position. Fig. 2 shows the macro
morphology of FSW joint and the fractured CT specimens after FCG test,
illustrating that the welded joint shows good quality with no obvious
weld defects.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image1/image1}

\end{center}
\end{figure}

Fig. 1 Schematic diagrams showing the sample orientations\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image2/image2}

\end{center}
\end{figure}

Fig. 2 Macro morphology of FSW joint and the fractured CT specimens

\textbf{3. Results}

3.1. Hardness and tensile properties tests

Fig. 3 shows the Vickers hardness arrangement in cross section of BM and
FSWed joint. It reveals that an entire drop of hardness occurred after
FSW as compared with BM. A roughly symmetrical micro-hardness
distribution could be seen in welded joint, and two well pronounced
softening bands with a distance of 10 mm from the joint center located
in AS-HAZ and RS-HAZ, respectively. The softening band in AS-HAZ was
about 5 mm in breadth, which was wider than the other one in RS-HAZ (3
mm). The average values of hardness in BM, AS-HAZ, RS-HAZ and WNZ were
186 HV, 124 HV, 132 HV, and 158 HV, respectively.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image3/image3}

\end{center}
\end{figure}

Fig. 3 Map of microhardness distribution along the transverse
cross-section of BM and welding joint.

The tensile properties of BM, AS-HAZ and WNZ were tested and shown in
Fig. 4, the results were consistent with hardness arrangement. BM showed
the maximum yield strength (YS: 515 MPa), ultimate tensile strength
(UTS: 593 MPa), while the YS and UTS in the softening position of AS-HAZ
were the lowest, decreasing to 262 MPa, 415 MPa respectively. In
contrast, WNZ processed the best elongation to failure and the strength
was lower than BM but much higher than AS-HAZ.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image4/image4}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image5/image5}

\end{center}
\end{figure}

Fig. 4 The engineering stress-strain curves and detail values of YS, UTS
and elongation to failure of BM, AS-HAZ and WNZ.

3.2. OM, EBSD and TEM observation

Fig. 5 shows the OM structure of welded joint. A homogeneous
microstructure formed in WNZ and less obvious ``S-line'' defects can be
observed (arrowed in the figure). A distinction among WNZ and HAZ can be
seen from the significant variation of grain size. Both AS-HAZ and
RS-HAZ maintained the similar pancake-shape grain structure as BM,
indicating that FSW heating input had little or no influence on the
grain structure in HAZ. Between WNZ and HAZ, there was a transition
zone, thermo-mechanically affected zone (TMAZ) characterized by a highly
deformed rheological structure due to the effects of shoulder-driven
flow and pin-driven flow. A clear boundary between TMAZ and HAZ can be
found in AS of joint, while the boundary in RS was not clear. According
to the microhardness arrangement map shown in Fig. 3, the softening band
with minimum hardness value did not located in TMAZ but in AS-HAZ.
Therefore, in current study, attention was focused on BM, WNZ and
softening band in AS-HAZ.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image6/image6}

\end{center}
\end{figure}

Fig. 5 OM image of FSW joint on cross section

The detail statistical information of grain structure of BM, WNZ and
softening band in AS-HAZ were further investigated by EBSD. The detailed
statistical fractions of recrystallized, sub-structured, deformed
structure were calculated and listed in the corresponding figures.
According to the EBSD maps shown in Fig. 6, it can be found that BM and
AS-HAZ mainly consisted of grains elongated along the rolling direction,
many sub-grains inside grains as well as some small recrystallized
grains in grain boundaries were distinguishable (Fig. 6a and 6c). In
WNZ, the grain size dramatically decreased and the intragranular
sub-grains entirely vanished (Fig. 6e), presenting a complete
recrystallized equiaxed-grained structure (Fig. 6f).\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image7/image7}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image8/image8}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image9/image9}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image10/image10}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image11/image11}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image12/image12}

\end{center}
\end{figure}

Fig. 6 EBSD misorientation color maps (a, c, e) and corresponding
recrystallization results (b, d, f): (a, b) BM, (c, d) AS-HAZ, (e, f)
WNZ

According to statistical information estimated by EBSD, the average
grain size (d*) and misorientation angle were further accurately
measured and displayed in Fig. 7. As can be seen that the d* of BM,
AS-HAZ were very close, with the values of 14 \selectlanguage{greek}μ\selectlanguage{english}m and 16 \selectlanguage{greek}μ\selectlanguage{english}m, while the d*
of WNZ was about 4 \selectlanguage{greek}μ\selectlanguage{english}m (Fig. 7a). In addition, both BM and AS-HAZ
processed high fraction of LAGBs structure, whereas, WNZ mainly
consisted of HAGBs recrystallized structure (Fig. 7b).\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image13/image13}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image14/image14}

\end{center}
\end{figure}

Fig. 7 Grain size (a) and misorientation distributions (b) of BM, AS-HAZ
and WNZ samples.

The characteristics of precipitates in BM, AS-HAZ and WNZ were
investigated by TEM and HRTEM viewed along with the
\textless{}110\textgreater{}\textsubscript{Al} projection. The results
are shown in Fig. 8, Fig. 9 and Fig. 10, respectively. In Fig. 8a, some
sub-structures remained in BM after solid solution treatment. Many
coherent Al\textsubscript{3}(Sc,Zr) particles with typical Ashby-Brown
strain characteristics can be observed in grains, however, the grain
boundary phase and precipitate-free zone (PFZ) were inconspicuous. At
higher magnification, many smaller short rod-shaped precipitates could
be distinguished clearly in the grains (Fig. 8b). According to the HRTEM
images and corresponding inverse fast Fourier transform (FFT) images,
one kind of these rod-shaped precipitates precipitated along
(11-1)\textsubscript{Al} and remained semi-coherent relationship with
matrix (framed by rectangle 1), were well accepted as ?' phases {[}25,
26{]}. Another kind of smaller precipitates consisted of several atomic
layers in thickness and width, and showed fully coherent with the matrix
(framed by rectangle 2 and 3), were known as GPII zones which often
found in Al-Zn-Mg-(Cu) alloys {[}27, 28{]}. TEM and HRTEM investigations
confirmed that GPII zones and ?' phases were the dominant precipitates
in BM.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image15/image15}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image16/image16}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image17/image17}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image18/image18}

\end{center}
\end{figure}

Fig. 8 TEM (a, b) and HRTEM images (c) of matrix precipitates in BM
sample. The inverse FFT images after mask to the phases in (c) image are
shown in (d).

As shown in Fig. 9a, sub-grains could also be observed in AS-HAZ. Some
discontinuous chain-typed ? (MgZn\textsubscript{2}) equilibrium phases
formed in grain boundary and the PFZ were obviously observed. As
compared with BM, the size of ?' precipitates in grains increased
slightly (Fig. 9b). It should be noted that the projections of some
precipitates showed short rod-shaped morphology while the others
displayed the round morphology, depending on the direction of
observation {[}27{]}. According to HRTEM observation, all of the ?'
precipitates remain semi-coherent relationship with matrix (Fig. 9c and
9d). No clear GPII zones were identified in AS-HAZ.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image19/image19}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image20/image20}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image21/image21}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image22/image22}

\end{center}
\end{figure}

Fig. 9 TEM images (a, b) and HRTEM image (c, d) of matrix precipitates
in the AS-HAZ sample.

In WNZ, sub-grains turned to be recrystallized structure due to high
weld heat input, and some Al\textsubscript{3}(Sc,Zr) particles grew
slightly (Fig. 10a). It is noted that the Ashby-Brown contrasts of the
grown Al\textsubscript{3}(Sc,Zr) particles become irregular, indicating
that the coherency would be break down during welding. It can be
observed that new ? equilibrium phase tended to attach to the grown
incoherent Al\textsubscript{3}(Sc,Zr) particles and overgrew. In
addition, a considerably amount of nanoscale rod-shaped ?' phase and
GPII zones also be observed at higher magnification (Fig. 10b).\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image23/image23}

\end{center}
\end{figure}\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image24/image24}

\end{center}
\end{figure}

Fig. 10 TEM images of matrix precipitates for the WNZ sample.

3.3. FCG behavior

The FCG rate of BM, AS-HAZ and WNZ are shown in Fig. 11. Generally, the
curve of FCG consists of three regions, namely, slow crack growth region
(I) in low stress intensity factor (\selectlanguage{greek}Δ\selectlanguage{english}K) range, stable propagation region
or Paris regime (II) in medium level range of \selectlanguage{greek}Δ\selectlanguage{english}K and rapid propagation
region (III) in high range of \selectlanguage{greek}Δ\selectlanguage{english}K. In present work, the investigated
range of \selectlanguage{greek}Δ\selectlanguage{english}K was about 7 to 30 MPa m\textsuperscript{1/2}. From 7 to 20
MPa m\textsuperscript{1/2}, all samples showed approximate linear
dependency of \selectlanguage{greek}Δ\selectlanguage{english}K on the FCG rate, indicated characteristic of Paris
regime (II). In this region, BM sample presented a consistently higher
FCG rate as compared with AS-HAZ and WNZ. The FCG rate of AS-HAZ rose
faster than WNZ with increasing of \selectlanguage{greek}Δ\selectlanguage{english}K. When \selectlanguage{greek}Δ\selectlanguage{english}K exceeded 20 MPa
m\textsuperscript{1/2}, a noticeable change in the slope of the curve
was observed, indicating that crack growth was entering into the rapid
propagation region (III). In this region, the FCG rate of BM and AS-HAZ
was close, much higher than that of WNZ.\selectlanguage{english}
\begin{figure}[H]
\begin{center}
\includegraphics[width=0.70\columnwidth]{figures/image25/image25}

\end{center}
\end{figure}

Fig. 11 Relationship of FCG rate and the stress intensity factor range
of BM, AS-HAZ and WNZ.

3.4. Morphological features of fracture surface Fig. 12 shows the
typical SEM fractographs of fatigue fracture features at stable
propagation (\selectlanguage{greek}Δ\selectlanguage{english}K = 15 MPa m\textsuperscript{1/2}) region. From the
fracture surfaces of BM and AS-HAZ, it can be seen that obvious tearing
ridges distributed along the propagation direction, tearing dimples and
small-scale facets formed during crack propagation. Some coarse
intermediate Fe-containing particles which usually formed during casting
were found in tearing dimples (Fig. 12a). At higher magnification,
fatigue striations on the facets can be observed clearly. The
characteristic fatigue fracture features of BM and AS-HAZ were
substantially similar but with two notable differences. Firstly, the
amount of tearing ridge on the fracture surface of AS-HAZ was fewer than
BM while the tearing dimples were more numerous. Secondly, the average
spacing of fatigue striations in AS-HAZ was smaller than BM, the
statistics average values were 1.5 \selectlanguage{greek}μ\selectlanguage{english}m and 1 \selectlanguage{greek}μ\selectlanguage{english}m, respectively. By
contrast, the fatigue fracture features of WNZ were quite different. The
tearing ridges turned to be smaller and not distributed strictly
parallel to the direction of propagation, and the amount of shallow
tearing dimples increased significantly. Large size intermediate
particles were probably broken down by pin tool during welding,
therefore coarse Fe-containing phases were scarcely observed in WNZ
(Fig. 12e). It is noted that many second cracks and facets with smaller
scale formed on the fracture surface of WNZ, the spacing of fatigue
striation became indistinguishable attributed to the existence of second
cracks. These fracture features revealed that the crack path of WNZ
would be more tortuous and in favor to slower FCG rate as compared with
BM and AS-HAZ. Fig. 12 SEM micrograph of fatigue fracture surfaces at
the stage of stable expanding region for BM, AS-HAZ and WNZ specimens.

\textbf{4. Discussion}

4.1. Effect of Sc on microstructure and mechanical properties

Significant recrystallization suppression effect of Sc addition in
aluminium alloys has been previously shown to arise from the
precipitation of second Al\textsubscript{3}(Sc,Zr) particles formed
during homogenization thermal treatment {[}16, 29{]}. In 7085Sc alloy,
Al\textsubscript{3}(Sc,Zr) particles formed in the matrix and inhibited
recrystallization during hot-rolled deformation as well as solution
treatment. Therefore, noticeable sub-grains remained in BM and produced
strengthening (Fig. 6 and Fig. 8a). Owing to sub-grains strengthening
and precipitation strengthening of GPII zones as well as ?' phases, the
UTS of 7085Sc-T6 increased to 593 MPa which was higher than AA7085-T6
reported in Ref. {[}30{]}.

The welding behaviors of 7085Sc alloy was also affected by
Al\textsubscript{3}(Sc,Zr) particles. It is known that grain growth and
dissolution/coarsening of hardening phases tend to occur when suffered
high FSW heat input, resulting significant softening in welded joint of
heat-treatable strengthening Al alloys {[}24, 31{]}. For the studied
7085Sc FSWed joint, sub-grains in AS-HAZ did not grew during FSW because
the migration of grain boundaries were retarded by
Al\textsubscript{3}(Sc,Zr) particles (Fig. 6 and Fig. 7). Therefore, the
presence of Al\textsubscript{3}(Sc,Zr) particles had significant benefit
to decrease the strengthening degradation in HAZ. The softening bands in
HAZ of 7085Sc FSWed joint could be mainly attributed to dissolution of
coherent GPII and growth of semi-coherent metastable ?' phases (Fig. 9),
similar to the previous reports {[}32, 33{]}.

As compared with HAZ, WNZ showed higher hardness and better tensile
properties (Fig. 3 and Fig. 4). One reason would be the high
strengthening of finer grain structure. It is known that high friction
heat input during FSW process would induce recrystallization and grain
growth in WNZ. For the 7085Sc alloy, grain growth was restrained in WNZ
due to high recrystallization resistance of the thermally stable
Al\textsubscript{3}(Sc,Zr) particles (Fig. 7). Another main reason would
be the strengthening of reprecipitated hardening phases. It has
estimated that the temperature in WNZ would be above 500\selectlanguage{ngerman}°C {[}34{]},
redissolution of precipitated phases would occur in this condition and a
supersaturated solid solution would form in WNZ during subsequent
air-cooling process. Similar to the slow quenching process of
Al-Zn-Mg-Cu alloys {[}35{]}, the supersaturated solid solution in WNZ
tended to decompose after welding, forming overgrowth quench-induced
equilibrium phases and nanoscale precipitates. It has been proved that \selectlanguage{english}?
equilibrium phases as the result of heterogeneous precipitation appear
to be mostly nucleated on pre-existing incoherent Cr or Zr-containing
dispersoids in slow quench processing of Al-Zn-Mg-Cu alloys {[}21,
36{]}. For 7085Sc alloy, Zr-containing dispersoids in the matrix were
mainly Al\textsubscript{3}(Sc,Zr), thus, the decomposition process of
the supersaturated solid solution in WNZ would be associated with
Al\textsubscript{3}(Sc,Zr) particles. As TEM observation showed that
just a few Al\textsubscript{3}(Sc,Zr) particles coarsened and lost
coherency when suffered the complicated thermo-mechanical input of FSW,
more coherent Al\textsubscript{3}(Sc,Zr) particles retained in WNZ after
welding (Fig. 10). Because coherent particles were able to reduce
heterogeneous precipitation of overgrowth \selectlanguage{english}? phases by preventing the
spread of voids {[}37{]}, only a small amount of incoherent
Al\textsubscript{3}(Sc,Zr) particles acted as nucleate of \selectlanguage{english}? equilibrium
phases. Therefore, the formation of \selectlanguage{english}? equilibrium phases in WNZ was
restrained, instead, many Zn and Mg solutes in the solid solution tended
to reprecipitate in form of \selectlanguage{english}?\selectlanguage{english}' and GPII as observed in Ref. {[}38{]}.
Consequently, mass of homogeneous fine hardening precipitates could be
obtained in WNZ, resulting improvement of hardness, YS and UTS values.

4.2. FCG behavior of welding joint

In current study, it was found that significant inhomogeneity of FCG
rate existed in different regions of 7085Sc FSWed joint (Fig. 11). As
widely known that FCG behavior is significantly associated with the
microstructure of alloys. For heat-treatable alloys, previous works have
proved that grain structure and hardening phases are the main
microstructure factors on FCG rate {[}31, 39, 40{]}. The BESD analysis
and TEM observation shown in Fig. 6 and Fig. 7 revealed that the grain
structures in BM and AS-HAZ were approximately similar, while the
morphology of hardening phases in these two regions were quite different
(Fig. 8 and Fig. 9). Thus, it can be reasonably concluded that the
difference of FCG behavior between BM and AS-HAZ was mainly induced by
hardening phases, while the effect of grain structure could be
negligible. As compared with BM/AS-HAZ, grain structure and hardening
phases in WNZ had more distinctive characteristic, therefore, the
effects of both two factors on the FCG behavior should be highlighted.
According to early researches {[}25, 31, 41, 42{]}, hardening phases
played a noticeable role to intervene FCG process by interaction between
dislocations. During the FCG process, dislocation nucleation and
movement were promoted at the crack tip, undoubtedly, the interaction
between hardening phase and movable dislocation turn to be frequent in
this region. For coherent or semi-coherent phase, dislocation could
shear it and glide reversibly with in a persistent slip band during
cyclic loading. This kind of kinematical reversibility of cyclic slip
would lead to a more deflective, tortuous and bifurcate fatigue crack
path as well as inhomogeneous deformation on the crack tip, contributing
to enhancement of FCG resistance. On the contrary, dislocation tends to
bypass non-coherent phases rather than shear. In this situation,
irreversible deformation turns to be dominated and increased damage
accumulation on the crack tip, and finally degrade the FCG resistance.
It should be noted that whether in shear or bypass mechanism, the radius
(\emph{r}) and the volume fraction (\emph{f}) of hardening phases are
two important microstructural parameters. Wen et al. {[}25{]} summarized
and proposed succinct equations to express the relationship as listed
below: shear mechanism:\(\frac{\text{da}}{\text{dN}}\propto\left(r\bullet f\right)^{-1/2}\)(1) bypass
mechanism:\(\frac{\text{da}}{\text{dN}}\propto{r\bullet f}^{-1/2}\) (2) Obviously, the FCG rate is negatively
related to \emph{r} and \emph{f} of shearable precipitates according to
Eq. (1), and the FCG rate is positively related to \emph{r} of
non-shearable precipitates while negatively related to its
\emph{f}according to Eq. (2). These relationships can be used to explain
the FCG inhomogeneity of 7085Sc FSWed joint. As can be seen from TEM
images, the dominated hardening phases in both BM and AS-HAZ samples
consisted mostly of coherent GPII and semi-coherent \selectlanguage{english}?\selectlanguage{english}', indicating that
shear mechanism was the main interaction mode in these two regions
during cyclic loading. However, the average size and volume fraction of
precipitates in AS-HAZ are much higher than that of BM. According to Eq.
(1), the FCG resistant of AS-HAZ should theoretically be higher than BM.
This theoretical deduction is entirely consistent with experimental data
shown in Fig. 11. For WNZ sample, precipitates in matrix consisted of a
considerably amount of reprecipitated coherent phase (\selectlanguage{english}?\selectlanguage{english}' and GPII) as
well as some coarse non-coherent phase (\selectlanguage{english}?). According to shear mechanism
mentioned above, high volume fraction of coherent phases would be
beneficial to enhancement of FCG resistance. The supporting evidences
could be found in Bray's work demonstrating that naturally aged Al-Cu-Mg
alloys had excellent fatigue properties {[}41{]}. According to Eq. (2),
bypass mechanism of \selectlanguage{english}? phases with large \emph{r} and low \emph{f} in WNZ
also contributed to reduction of FCG rate. In addition to the effects of
precipitates, another considerable factor on FCG rate of WNZ is the
grain structure. It has been known that due to small angle of
preferential slip planes between adjacent grains, fatigue crack tends to
expand across multiple grains at a time in alloy with high texture
intensities, resulting low FCG resistance {[}43, 44{]}. As compared with
BM and AS-HAZ, WNZ possessed recrystallized grains with smaller size and
random crystal orientation (Fig. 6 and 7), in this situation, a higher
force was required for the propagation of fatigue crack as the impeding
effect caused by high angle grain boundary. Therefore, the deflection,
torsion and bifurcation of fatigue crack occurred more frequently (Fig.
12), leading to a significant decrease of FCG rate. Due to the combined
effects of precipitates and grain structures, FCG property of WNZ was
better than other welding regions as well as BM. In addition, the
influence of grains structure on FCG is associated with the cyclic
plastic zone (CPZ) at the crack tip {[}25{]}. The CPZ size could be
estimated by\(\delta_{f}=\frac{1}{8\pi}\left(\frac{K}{\sigma_{y}}\right)^{2}\), here (\selectlanguage{greek}σ\selectlanguage{english}\textsubscript{y}) accurately
would be the strength of the plastic zones near the fatigue crack tip.
Obviously, the CPZ rises with increasing of \selectlanguage{greek}Δ\selectlanguage{english}K. When CPZ size is larger
than a grain scale, the effects of grain structure on the FCG would turn
to be dominant {[}45{]}. Because BM and AS-HAZ processed similar grain
structure, thus, the crack in BM and AS-HAZ samples showed roughly the
same FCG rate at high \selectlanguage{greek}Δ\selectlanguage{english}K (\textgreater{}20 MPa m\textsuperscript{1/2}).

\textbf{5. Conclusions}

The inhomogeneity of microstructure, hardness, tensile properties and
FCG behavior of 7085Sc FSW weld joint was investigated and discussed.
Based on the experimental results, the following major conclusions can
be drawn:

(1) The recrystallization resistance of the 7085Sc alloy was improved by
coherent Al\textsubscript{3}(Sc,Zr) particles, thus, sub-grains were
remained in BM after solution treatment and the grain growth was
inhibited in AS-HAZ during FSW process. Frictional heating and intense
plastic deformation caused fined recrystallized grains formed in WNZ,
the average grain size was 4 \selectlanguage{greek}μ\selectlanguage{english}m, smaller than that of BM (14 \selectlanguage{greek}μ\selectlanguage{english}m) and
AS-HAZ (16 \selectlanguage{greek}μ\selectlanguage{english}m).

(2) FSW heat input induced dissolution of GPII and slightly growth of
metastable phases ?' in AS-HAZ, by contrast, completely redissolution of
hardening precipitates occurred in WNZ forming supersaturated solid
solution which tend to decompose to form new reprecipitated GPII, ?' and
? phases after FSW. However, the heterogeneous precipitation of
overgrowth ? phases were reduced by coherent Al\textsubscript{3}(Sc,Zr)
particles. Owing to precipitation strengthening and fine grain
strengthening, the tensile properties of WNZ (UTS: 330MPa and YS:
485MPa) were higher than that of AS-HAZ (UTS: 415MPa and YS:262MPa).

(3) The FCG rate of AS-HAZ was lower than BM in Paris regime because the
FCG resistance in HAZ was improved by slight grew semi-coherent ?' as
shear mechanism. However, the FCG rates of BM and AS-HAZ were close at
high \selectlanguage{greek}Δ\selectlanguage{english}K due to the similar grain structures. WNZ possessed the best
fatigue endurance because of shear mechanism of reprecipitated coherent
phases (GPII and ?') and bypass mechanism of ? phases with large size
and low volume fraction. Moreover, fine recrystallized grains with
random crystal orientation in WNZ also contributed to reduction of FCG
rate. The combined effects of precipitates and grain structures induced
frequent deflection, torsion and bifurcation of fatigue crack, resulting
in significant reduction of FCG rate in WNZ.

\textbf{CRediT authorship contribution statement}

\textbf{Zixing Jia:} Methodology, Investigation, Writing - original
draft. \textbf{Lili Wei:} Funding acquisition, Writing - review \&
editing. \textbf{Hongfeng Huang:} Funding acquisition,
Validation.\textbf{Chongyu Liu:} Conceptualization, Supervision.
\textbf{Hongjie Jiang:} Methodology, Investigation. \textbf{Xuda Xu:}
Resources.

\textbf{Acknowledgments}

This work was funded by the Natural Science Foundation of Guangxi (grant
no.2018GXNSFAA138065, 2018GXNSFBA281074) and the Science and Technology
Project of Guangxi (grant no. GKAD19110041, GKAD19110087).

\textbf{Data Availability Statement}

All data included in this study are available upon request by contact
with the corresponding authors.

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