479
18
Heat-Affected-Zone Degradation of Mechanical Properties
This chapter deals with the degradation of mechanical properties in the heat-affected zone (HAZ) caused by solidstate transformations, such as grain coarsening, recrystallization and grain growth in work-hardened materials,
overaging in Al alloys, dissolution of strengthening precipitates in heat-treatable Ni-base alloys, and tempering of
martensite in dual-phase steels.
18.1 ­Grain Coarsening
Figure 18.1 shows grain coarsening in the HAZ of a 430 ferritic stainless steel weld made by gas−tungsten arc
welding (GTAW) [1]. Grain coarsening is also evident in the HAZ of a 312 duplex stainless steel weld made by
GTAW, as shown in Figure 18.2 [2]. In both cases grain coarsening in the HAZ causes coarse grains to grow at the
fusion line by epitaxial growth and grow as coarse columnar grains into the bulk fusion zone.
Figure 18.3 shows the thermal cycles in the HAZ during welding. The closer to the fusion line, the higher the
peak temperature is and the longer the material stays at high temperatures to increase diffusion and hence grain
growth. In electroslag welding of steels, the cooling rates are very slow and the HAZ can stay at high temperatures
for a long time. Consequently, severe grain growth can occur in the HAZ and hence the fusion zone (Figure 1.25),
and the toughness of the weld can be very poor. Coarse grains reduce not only strength but also toughness. In
narrow-gap electroslag welding [3], the gap between the two plates to be butt welded is reduced, and so is the heat
input that is needed to maintain a smaller weld pool and slag pool during welding. This lower heat input increases
the cooling rate, which reduces grain growth and improves the weld toughness.
Fine precipitates of carbides, nitrides, or carbonitrides can help inhibit grain growth in the HAZ of steels [4].
Precipitates such as titanium nitride TiN that resist coarsening and dissolution better at elevated temperatures, i.e.
more stable, can be more effective in inhibiting grain growth. They can remain as fine dispersed precipitates to pin
down grain boundaries and inhibit grain growth.
Wadsworth et al. [5] showed severe grain growth in the HAZ of a molybdenum (Mo) weld made by electronbeam welding. Similar grain growth has been observed in a vanadium (V) weld made by laser-beam welding [6].
Significant grain growth still occurs in the HAZ in spite of the relatively low heat input and high cooling rate in
electron- or laser-beam welding. The melting points of Mo (2623 °C) and V (1910 °C) are very high, which increase
the peak temperature of the thermal cycle and the high-temperature residence time. With arc welding, grain
growth in these materials can be even worse.
Tabernig and Reheis [7] tensile tested Mo before and after electron beam welding (EBW) and observed significant reductions in the strength and ductility by welding. The severe loss of ductility suggests a severe loss of
toughness after welding.
Welding Metallurgy, Third Edition. Sindo Kou.
© 2021 John Wiley & Sons, Inc. Published 2021 by John Wiley & Sons, Inc.
480
18 Heat-Affected-Zone Degradation of Mechanical Properties
Figure 18.1 Grain growth near fusion line of weld of 430 ferritic stainless steel (~Fe-16Cr) made by GTAW: (a) photo of top
of weld; (b) enlarged. Source: Yang and Kou [1].
Figure 18.2 Grain growth near fusion line of weld of 312 duplex stainless steel (~Fe-30Cr-9Ni) made by GTAW. Transverse
cross-section of weld. Source: Yu and Kou [2].
Figure 18.3 Grain growth in HAZ: (a) transverse cross-section of weld and thermal cycles; (b) grain growth.
18.2 ­Recrystallization and Grain Growth
Recrystallization and grain growth can occur when welding a work-hardened material. Figure 18.4 shows schematically the weakening in the HAZ due to recrystallization and grain growth during welding. For a given material, hardness has been shown to be proportional to the yield strength [8]. Since the HAZ of a weld is usually too
18.2 ­Recrystallization and Grain Growt
Figure 18.4 Effect of welding on work-hardened material: (a) before welding; (b) after welding. Hardness has been shown
to be proportional to the yield strength.
Figure 18.5 Microstructure of the weld of a work hardened 304 stainless steel: (a) original base metal; (b) base metal after
work hardening; (c) grain-boundary carbide precipitation (sensitization at 600–850 °C); (d) recrystallization; (e) grain
growth; (f) columnar dendrites. Source: Metals Handbook [9]. © ASM International.
small for tensile testing, microhardness measurement across the HAZ of a material is a convenient way to study
the effect of welding on the material.
Evidence of recrystallization and grain growth can be seen in Figure 18.5, which shows the transverse micrographs (micrographs taken from the transverse cross-section) of the HAZ of a work-hardened 304 stainless steel
[9]. At position d the grains are much smaller than those in the base metal (BM) at position b, where the slip bands
in the grains indicate severe deformation of the grains caused by work hardening of the workpiece before welding.
These small grains indicate recrystallization caused by heating during welding. On the contrary, near the fusion
line at position e the grains are significantly larger than those in the base metal at position b. These coarse grains
indicate grain growth caused by heating during welding. HAZ grain growth is also common in other work-hardened materials, e.g. in GTAW of AZ31-H24 Mg alloy sheets [10].
Figure 18.6a shows the microhardness profiles of two arc welds of 5083 Al alloy (~Al-4.0Mg) [11], one workhardened before welding to a higher hardness level than the other. However, it also suffers a greater loss of
strength after welding. The resultant strength in the HAZ (and fusion zone) appears to be close.
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18 Heat-Affected-Zone Degradation of Mechanical Properties
Figure 18.6 Hardness profiles across welds of work-hardened materials: (a) arc welded 5083 Al of two hardened levels.
Cook, Shannon, and Hard [11], Welding Journal, 1955, © American Welding Society; (b) friction-stir-welded Cu-37Zn brass.
Source: Cam et al. [12]. Welding Journal, November 2009, © American Welding Society.
Figure 18.6b shows the microhardness profiles across a weld of Cu-37Zn made by friction-stir-welding at the rotation speed of 1600 rpm and the travel speed of 175 mm/min [12]. The workpiece was work-hardened before welding. Significant loss of strength (as reflected by hardness) is evident in the HAZ. The HAZ was annealed by the
friction heat generated by tool rotation during welding. The smaller grains in the stir zone (caused by dynamic
recrystallization induce by tool rotation) could have contributed to the higher strength in the stir zone (near centerline) than the HAZ. According to the Hall-Petch equation, the yield strength σy increases with decreasing grain
size d as follows:
y
a
b
d
(18.1)
where a and b are constants related to properties of materials [13].
Figure 18.7 shows that increasing heat input can significantly increase the HAZ width and reduce the HAZ strength
of work-hardened materials after welding. As shown by Eq. (2.17), increasing the heat input decreases the cooling rate.
This tends to broaden the thermal cycle and prolong the high-temperature residence time for diffusion and hence
Temperature T,
482
Figure 18.7 Effect of heat input on welding of work-hardened materials: (a) (b) lower heat input; (c) (d) higher heat input.
18.3 ­Overaging in Al Alloy
Figure 18.8 Effect of heat input on HAZ hardness in arc welding of
5356 Al (~Al-5Mg). Source: White et al. [14], Welding Journal, November
1960, © American Welding Society.
recrystallization and growth and result in a wider and weaker HAZ. Figure 18.8 shows the HAZ strength decreases and
the HAZ width increases and with increasing heat input in arc welding of 5356 Al [14]. Lee et al. [15] also shows that,
in friction stir welding (FSW) of AZ31-H24 Mg, the grain size in the stir zone increases with increasing tool rotation
speed and decreasing tool travel speed. The joint strength of work-hardened Mo sheets welded by EBW was higher than
that by GTAW [16]. In EBW of 0.2 mm Mo sheets, the HAZ was narrow though the loss of strength was significant. In
GTAW of 0.4 mm Mo sheets, however, the strength was reduced over a much wider HAZ.
18.3 ­Overaging in Al Alloys
18.3.1 Al-Cu-Mg (2000-Series) Alloys
18.3.1.1
Microstructure and Strength
2219 Al is close to a binary alloy of Al-6.3Cu and thus can be used as model for studying the response of 2000-series
Al alloys to welding. Figure 18.9 shows the TEM (transmission electron microscopy) images in the HAZ of a 2219
Al alloy that has been pre-heat-treated to contain θ′ precipitates (Figure 17.5a) before welding [18]. The volume
Figure 18.9 Microstructure in a 2219 Al that has been aged to contain θ′ precipitates before welding. Source: Dumolt et al.
[17]. © Welding Research Council.
483
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18 Heat-Affected-Zone Degradation of Mechanical Properties
fraction of θ′ decreases from the base metal toward the fusion boundary. The peak temperature during welding
increases from the base metal toward the fusion line, causing first partial and then full θ′ reversion as the fusion
line is approached. On the average, the θ′ in the middle image is fewer in number and smaller in size though
larger θ′ can also be seen. The larger θ′ appears to grow at the expense of the smaller one, that is, θ′ reversion is
accompanied by θ′ coarsening. It is worth mentioning that 2219-T6 Al has been shown to contain θ′ as the primary precipitation strengthening phase [19–21]. Thus, Figure 18.9 is relevant to welding of 2219-T6 Al.
The microstructure in the 2219 Al weld in Figure 18.9 can be explained with the help of Figure 18.10. Points 1, 2,
and 3 are inside the HAZ. Point 3 has the lowest peak temperature, and some θ′ reversion occurs, that is, partial reversion. At point 2, the peak temperature is higher, and more reversion occurs. Coarsening of some θ′ is evident, suggesting overaging has occurred. Point 1 is closest to the fusion line, and its highest peak temperature causes full reversion
of θ′. The cooling rate here is too high for θ′ to reprecipitate. This results in a supersaturated solid solution at point 1.
The reversion of θ′ weakens the HAZ as shown schematically by the microhardness profile in the as-welded (AW)
condition. At point 1, the supersaturated solid solution allows GP zones to form spontaneously at room temperature.
This can cause some hardness increase at point 1 after post-weld natural aging (PWNA). Overaging at point 2 keeps
the hardness from full recovery in post-weld artificial aging (PWAA). If it is possible to solution-heat-treat the entire
workpiece after welding, quench it and then artificially age it, the full strength can be recovered in the HAZ. However,
solution heat treatment can be difficult for a large workpiece and quenching can cause distortion.
Figure 18.11 shows the TEM images in the HAZ of a 2219 Al alloy that has been prepared to contain GP zones
before welding [17]. The image on the right shows small particles of GP zones in the base metal. (The dark bands
are artifacts caused by the sample preparation procedure for TEM). The welding heat causes GP zones to dissolve
Figure 18.10 HAZ of an Al-Cu alloy prepared to contain θ′ (similar to T6 temper) before welding: (a) phase diagram; (b)
thermal cycles; (c) precipitation C-curve; (d) as-welded microstructure; (e) microhardness profiles in as-welded (AW),
postweld natural aging (PWNA) and postweld artificial aging (PWAA) conditions.
18.3 ­Overaging in Al Alloy
Figure 18.11 Microstructure in a
2219 Al that has been aged to
contain GP zones before welding.
Source: Dumolt et al. [17]. ©
Welding Research Council.
Figure 18.12 HAZ of an Al-Cu
alloy prepared to contain GP zones
(similar to T4 temper) before
welding: (a) phase diagram;
(b) thermal cycles; (c) θ′
precipitation C-curve; (d) as-welded
microstructure; (e) microhardness
profiles in as-welded (AW),
postweld natural aging (PWNA) and
postweld artificial aging (PWAA)
conditions. θ in base metal not
shown.
485
486
18 Heat-Affected-Zone Degradation of Mechanical Properties
in the HAZ, i.e., GP-zone reversion. However, it also causes θ′ to precipitate in region III of the HAZ as shown by
the TEM image on the left. Figure 18.11 is relevant to welding 2219 in the T4 condition.
The microstructure in the 2219 Al weld in Figure 18.11 can be explained with the help of Figure 18.12. The
C-curve shown in Figure 18.12c is a temperature-time curve that shows the time required for θ′ to precipitate as a
function of temperature. At point 2 the peak temperature during welding reaches the precipitation temperature
range of θ′ and causes it to precipitate. This can cause a small peak to form at point 2 in the as-welded (AW) condition as shown schematically in Figure 18.12e. During PWNA, GP zones can form from the supersaturated solid
solution everywhere in the HAZ except at point 2, where some θ′ already exists. Upon PWAA, the HAZ can
recover strength significantly by precipitating θ′. This recovery of strength, however, is less significant at point 2
because some θ′ already exists and some overaging can occur during PWAA.
The explanations for the response of Al-Cu alloys to welding presented in Figures 18.10 and 18.12 are consistent
with the results in the recent study of Lin et al. [22] on 2219 Al alloy. These explanations are also consistent with
the earlier results of Kou and Le [23] on 6061 Al alloy, which will be presented subsequently. Figure 18.13 shows
the transverse cross-section of 2024 Al lap-welded in the T3 temper [24], which is solution heat treated, cold
worked and naturally aged. T3 is similar to T4 except for cold working before natural aging. The sample was
etched with an aqueous solution containing 0.5% HF. This sample is similar to that in Figure 18.11 in the sense
that a solution region exists near the fusion zone and a precipitation region exists next to the solution range.
Lin et al. [22] welded 2219-T6 by variable-polarity GTAW. As shown in Figure 18.14. The microhardness
decreases significantly across the HAZ. PWAA does not increase the HAZ hardness because it is overaged during
welding. Since the Al-rich dendrites in the fusion zone are likely supersaturated solid solution, they can gain
strength significantly by PWAA. However, the Cu in the interdendritic areas still exists as Cu-rich eutectic, unable
to help increase the hardness by PWAA. The TEM image of the HAZ in the as-welded condition shows coarse
Figure 18.13 Transverse cross-section of 2024 Al (~Al-4.4Cu-1.5Mg) lap-welded in T3 temper (solution heat treated, cold
worked and naturally aged): (a) macrograph; (b)–(f) micrographs. Source: Soysal, Yu, Kou [24].
18.3 ­Overaging in Al Alloy
Figure 18.14 2219 Al (~Al-6.3Cu) welded in T6 condition by gas−tungsten arc welding: (a) microhardness profiles; (b) TEM
image in HAZ in as-welded condition showing coarse equilibrium phase θ (Al2Cu) as evidence of overaging occurring during
welding in T6 condition. Source: Lin et al. [22]. © Taylor and Francis.
Figure 18.15 2219 Al (~Al-6.3Cu) welded after solution heat treating plus quenching (similar to T4 temper, assuming GP
zones had formed before welding) by gas-tungsten arc welding: (a) microhardness profiles; (b) TEM image in HAZ in
as-welded condition showing metastable phase θ′ as sign of its precipitation during welding. Source: Lin et al. [22]. © Taylor
and Francis.
precipitates of θ (Al2Cu), which is clear evidence of overaging. These precipitates are much smaller than the
Al2Cu particles in the base metal of 2219 Al (Figure 15.6b).
The 2219-T6 Al was also welded by Lin et al. [22] after it had been solution heat treated at 535 °C for 90 minutes
and then quenched. The results are shown in Figure 18.15. The microhardness profile in the as-welded condition
differs from that in Figure 18.14 in two significant ways. First, the microhardness of the base metal is lower here
because of the solution heat treating before welding. Second, a microhardness peak exists in the HAZ. This peak
suggests θ′ precipitation here during welding, which is confirmed by the TEM image showing θ′ at the peak
(Figure 18.15c) but not elsewhere in the HAZ (Figure 18.15b). The peak here is similar to that in the HAZ shown
previously in Figure 18.12e, where the workpiece has been prepared to contain GP zones before welding. The 2219
Al here could have also contained GP zones sometime after solution heat treatment, but no higher-magnification
TEM images were taken to show the GP zones. Although a significant drop still exists in the HAZ after PWAA, it
is significantly less severe than the one shown in Figure 18.14a.
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18 Heat-Affected-Zone Degradation of Mechanical Properties
Figure 18.16 Effect of rotation speed on underwater friction stir welding of 2219-T6: (a) transverse cross-section of weld;
(b) θ′ in base metal (BM); (c) θ′ in HAZ at 50 mm/min travel speed; (d) θ′ in HAZ at 150 mm/min travel speed; (e)
microhardness profiles. Source: Liu et al. [21]. © Elsevier.
Chen et al. [19, 20] studied the precipitates in FSW of 2219-T6 Al. They showed evidence of overaging and next
to the stir zone based on both TEM images microhardness measurements taken across the weld.
In summary, welding a 2000-series Al alloy in the T6 condition can result in significant overaging and the loss of
hardness (strength) because this overaging cannot be recovered effectively by PWAA. If appropriate, welding in the
T4 condition is preferred. Overaging can be less, and the recovery of hardness by PWAA can be more effective.
18.3.1.2
Effect of Welding Parameters or Process
Liu et al. [25] conducted FSW of 2219 Al in cooling water and reduced the extent of overaging. Water cooling can
increase the cooling rate and reduce the high-temperature residence time to reduce overaging. They also studied
the effect of welding parameters in underwater FSW, including the effect of the rotation speed on the ­microstructure
and strength in underwater FSW of a 7.5 mm-thick 2219-T6 Al [21]. As shown in Figure 18.16, when the travel
speed is increased from 50 to 150 mm/min at 800 rpm, the θ′ in HAZ become finer in size and larger in number.
(A similar change was observed in the TMAZ.) The increase in travel speed causes the HAZ hardness to increase,
consistent with the TEM images. The hardness profiles and TEM images both indicate significant overaging when
welding in the T6 condition even underwater.
A welding process with a lower heat input can help reduce overaging in welding a 2000-series Al alloy. Figure 18.17
shows the results of welding a 4-mm-thick 2219-T87 Al [26]. T87 refers to solution that is heat treated, cold worked
by a thickness reduction of 10%, and then artificially aged. As compared to GTAW, FSW causes a narrower HAZ
and slightly less overaging because of its less heat input. Martukanitz and Howell [27] showed microhardness
18.3 ­Overaging in Al Alloy
Figure 18.17 Effect of welding process on microhardness profile: (a) gas−tungsten arc welding; (b) 2219-T87 Al by friction
stir welding; T87 involves solution-heat treating, cold working and then artificially aging. Source: Lei et al. [26],
© SpringerNature.
Figure 18.18 Effect of arc oscillation on HAZ width in GTAW of 2014-T6 Al (~Al-4.4Cu): (a) wider HAZ without oscillation;
(b) narrower HAZ with 1 Hz oscillation at 1.9 mm amplitude. 60A, 11 V 4.2 mm/s torch travel speed. Source: Kou and Le [28].
Welding Journal, March 1985, © American Welding Society.
profiles in the HAZs of 2195-T8 Al alloy, where T8 refers to solution heat treated, cold worked, and then artificially
aged. As compared to variable-polarity plasma arc welding (VPPAW), laser beam welding (LBW) produces a much
narrower HAZ and causes much less overaging. Figure 18.18 shows that transverse arc oscillation at a low frequency (e.g. 1 Hz) can reduce the width of the HAZ significantly [28]. As explained in Figure 8.18, such arc oscillation can increase the weld pool travel speed V significantly and hence significantly decrease Q/V.
18.3.2 Al-Mg-Si (6000-Series) Alloys
18.3.2.1
Microstructure and Strength
For Al-Mg-Si alloys the precipitation sequence, Eq. (17.5), is as follows:
Al-Mg-Si e.g., 6061 : SS
GP
Mg2Si
Mg2Si
(18.2)
489
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18 Heat-Affected-Zone Degradation of Mechanical Properties
Figure 18.19 Microhardness profiles in HAZs of 6061 Al: (a) welded in T6 temper; (b) welded in T4 temper. Source: From
Kou and Le [23].
The precipitation-hardening curves during artificial aging of 6061 Al (~Al-1Mg-0.6Si) have been shown in
Figure 17.6 [29].
Kou and Le welded 6061-T4 Al and 6061-T6 Al by GTAW [23]. As mentioned previously, the T6 temper refers to
solution heat treatment followed by artificial aging [30]. The microhardness profiles in Figure 18.19a for T6 are
similar to those shown schematically in Figure 18.10 for an Al-Cu alloy heat treated to contain θ′ before welding.
The similarity includes a sharp hardness decrease in the HAZ caused by β′ reversion during welding and a region
of very low hardness after PWAA, where overaging occurred during welding. This suggest the hardness profiles in
Figure 18.19a can be explained based on Figure 18.10.
The T4 temper refers to solution heat treatment followed by naturally aging to a substantially stable condition
[30]. Thus, a T4 aluminum workpiece can be expected to contain GP zones before welding. In Figure 18.19b for
T4 the two microhardness profiles for the as-welded condition and PWNA are similar to those shown schematically in Figure 18.12 for an Al-Cu alloy heat treated to contain GP zones before welding. The similarity includes a
mild hardness decrease in the HAZ caused by reversion of the GP zones and a small hardness peak in the HAZ
caused by β′ precipitation during welding. As for the microhardness profile for PWAA, the location of the minimum hardness is close to that of the small hardness peak. It is not clear why a small bump exists in the HAZ near
the base metal.
Figure 18.20 shows the microhardness profiles of 6061 Al welded in T4 and T6 temper after PWNA and PWAA
[31]. Under the conditions involved in welding and aging, HAZ weakening is significantly less if 6061 Al is welded
in the T4 temper instead of T6. This, in fact, can also be seen in the results of 2219-T6 Al shown in Figures 18.14
and 18.15.
Figure 18.21 shows the microhardness profiles in a 6063 Al (~Al-0.48Mg-0.44Si) welded by FSW in the T5
temper, which refers to extrusion, solution heat treatment at 550 °C, and artificially aging at 205 °C for 1 hour
[32]. The arrowheads indicate locations of the TEM images in Figure 18.22. The locations of the images are as
follows: BM for the region with same hardness as the base metal, LOW for the HAZ region lower in hardness
than the base metal, MIN for minimum hardness region in the HAZ, and SOF(0) the softening region in the
HAZ. Similar to the as-received base metal, the BM region consists of a high density of fine needle-shaped
precipitates (which were not identified) and a low density of much coarser β′ precipitates. The fine
­needle-shaped precipitates ­dissolve partially in the LOW region and completely in the MIN and SOF(0)
regions. In the LOW region the needle-shaped precipitates grow into β′ precipitates, that is, overaging. The
MIN region contains only a low density of β′ precipitates, similar to the BM region. The SOF(0) region has no
18.3 ­Overaging in Al Alloy
Figure 18.20 Welding 6061 Al in: (a) T6 temper; (b) T4 temper. Source: Redrawn from Metzger [31]. Welding Journal,
October 1967, © American Welding Society.
Figure 18.21 Microhardness profiles across friction stir weld along the mid-thickness of a 6 mm-thick 6063-T5 Al. T5
temper: extruded, solution heat treated at 550 °C, and aged artificially at 205 °C for 1 hour. Source: Sato et al. [32].
© SpringerNature.
precipitates of any kind. PWAA was conducted at 175 °C for various length of time. It significantly increased
the density of the fine needle-shaped precipitates and hence the hardness except for the LOW region, where
overaging has occurred.
It can be summarized from the above discussion that, as in the case of 2000-series Al alloys, welding a 6000-series
Al alloy in the T6 condition can result in significant overaging and PWAA cannot recover the loss of hardness
(strength) effectively. Welding in the T4 condition is preferred, as overaging can be less and the recovery of hardness by PWAA can be more effective.
18.3.2.2
Effect of Welding Processes and Parameters
Figure 18.23 shows the effect of the heat input per unit length of the weld on the HAZ microhardness profile of
6061-T4 [33]. The heat input per unit length of the weld is Q/V, where Q is the power input (e.g. I × E in arc welding) and V the travel speed. As shown in Eq. (2.17), decreasing Q/V increases the cooling rate and reduces the
high-temperature residence time to reduce overaging. Figure 18.24 shows the HAZ in 6061-T6 Al is much wider
in gas−tungsten arc (GTA or TIG) welding than in laser-beam welding [34].
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18 Heat-Affected-Zone Degradation of Mechanical Properties
Figure 18.22 TEM images of as-welded 6063-T5 Al at locations indicated in Figure 18.21. Source: Sato et al. [32].
© SpringerNature.
18.3.3 Al-Zn-Mg (7000-Series) Alloys
For Al-Zn-Mg alloys, the precipitation sequence is as follows:
Al-Zn-Mg e.g., 7005 : SS
GP
Zn 2 Mg
Zn 2 Mg
(18.3)
Al-Zn-Mg alloys tend to age significantly more slowly than Al-Cu-Mg (2000 series) and Al-Mg-Si alloys
(6000 series). For instance, the artificial aging of 7005 Al (~Al-4.5Zn-1.2Mg) is much slower than that of 2014
Al (~Al-4.5Cu-0.6Mg) [35]. Consequently, Al-Zn-Mg alloys has a smaller tendency to overage during welding
than Al-Cu-Mg and Al-Si-Mg alloys. Another interesting aspect of Al-Zn-Mg alloys is that they can recover
strength nearly fully after welding by natural aging. An example has been shown previously in Figure 17.7.
Figure 18.25 shows that 7005 Al alloy can almost fully recover its strength in the HAZ by natural aging
[23]. This can occur even when the alloy is welded after artificial aging. Figure 18.26 shows the microhardness
Figure 18.23 Less overaging in welding 6061-T4 Al with a smaller
heat input. Source: Burch [33]. Welding Journal, 1958, © American
Welding Society.
18.3 ­Overaging in Al Alloy
Effect of welding process/condition on HAZ of 6061-T6 AI
Vickers hardness (Hv; load 1.96N)
110
Figure 18.24 Effect of welding process and condition on
HAZ width of 6061-T6 Al. Source: Hirose et al. [34].
© SpringerNature.
100
90
80
70
Laser (133 mm/s)
Laser (66.7 mm/s)
TIG (5 mm/s)
60
50
40
0
2
4
6
8
10
Distance from fusion boundary (mm)
12
Figure 18.25 Microhardness profiles in HAZs of 7005 Al (~Al-4.5Zn-1.2Mg) alloy: (a) naturally aged before welding;
(b) artificially aged at 130 °C for 1 hour before welding. 1 : 3 hours; 2 : 4 days; 3 : 30 days; 4 : 90 days. Circles indicate peak
temperatures reached during welding. Source: Mizuno [36], © Japan Welding Society.
Figure 18.26 Microhardness profiles in HAZs of 7146 Al (~Al-7.1Zn-1.3Mg) alloy: (a) naturally aged before welding; (b)
annealed before welding. Source: From Kou and Le [23].
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18 Heat-Affected-Zone Degradation of Mechanical Properties
profiles in the HAZs of 7146 Al (~Al-7.1Zn-1.3Mg) [36]. When welded after natural aging, the HAZ hardness
recovers fully in eight days. When welded in the O-temper, the HAZ can age naturally to a significantly
higher hardness level than the base metal. As mentioned in Section 17.2, in the O-temper an alloy is annealed
to achieve the lowest-strength condition in order to maximize subsequent workability [30]. Kalemba et al.
[37] showed 7136-T76 Al (~Al-7.9Zn-2.0Mg-1.9Cu) can recover strength mildly by natural aging after FSW.
The recovery is not as impressive as that in Al-Zn-Mg alloys, probably due to some overaging effect caused
by alloying with Cu. Based on both Gleeble temperature simulation and TEM, Zhang et al. [38] showed the
effect of welding thermal cycles on the microstructural evolution in an Al alloy essentially Al-4.2Zn-1.0Mg
in composition.
18.4 ­Dissolution of Precipitates in Ni-Base Alloys
Figure 18.27 shows schematically the weakening of the HAZ of a heat treatable Ni-base alloy welded in the
aged condition. The thermal cycles show that the area next to the fusion line is heated above the precipitation
C-curve of γ′. This causes partial reversion near the base-metal side of the HAZ (point 2) and full reversion
near the fusion-line side (point 1). Reversion of the strengthening phase γ′ reduces the hardness or strength
in the HAZ.
Owczarski and Sullivan [39] welded a Udimet 700 alloy that had been fully aged by the following heat-treatment
procedure before welding: (i) 1165 °C for 4 hr + air cool (solution); (ii) 1080 °C for 4 hr + air cool (primary age); (iii)
Figure 18.27 Reversion of γ′ in HAZ: (a) phase diagram; (b) thermal cycles; (c) precipitation C-curve; (d) microstructure; (e)
hardness distribution.
18.4 ­Dissolution of Precipitates in Ni-Base Alloy
Figure 18.28 Microstructure of Udimet 700 weld: (a) as-received material; (b) initial solution (fine γ′ dissolved);
(c) further solution (coarse γ′ dissolving); (d) advanced stage of solution (coarse γ′ nearly all dissolved); (e) weld metal
containing fine γ′. Source: Owczarski et al. [39]. Welding Journal, September 1964, © American Welding Society.
840 °C for 4 hr + air cool (intermediate age); and (iv) 760 °C for 16 hr + air cool (final age). As mentioned in
Figure 17.8, the matrix γ is strengthened by the precipitates of γ′, i.e. Ni3(Al, Ti). As shown in Figure 18.28, the
base-metal microstructure consists of coarse angular γ′ and fine spherical γ′ between the coarse γ′. Reversion (dissolution) of γ′ occurs in the HAZ, with disappearing of fine γ′ and rounding of coarse angular γ′ on the base-metal
side of the HAZ, partial reversion of coarse γ′ in the middle of the HAZ, and complete reversion of coarse γ′ on
the fusion-zone side of the HAZ. With much Ti and Al released by reversion, the areas near the reverted coarse γ′
become supersaturated with Ti and Al. This supersaturation causes finer γ′ to re-precipitate during cooling (Figure
18.28c). The weld metal contains a uniform distribution of fine γ′ precipitate.
Jun et al. [40] showed the dissolution of primary γ′ precipitates (spherical) in a Waspaloy caused by FSW. The
alloy was Ni-19.14Cr-13.11Co-4.12Mo-3.13Ti-1.31Al-0.76Fe-0.053C in wt%, solutionized and then double aged at
843 °C for 4 hours and 732 °C for 4 hours. In the base metal, coarse spherical precipitates of primary γ′ were visible, but fine precipitates of secondary γ′ were also present. Partial dissolution of coarse γ′ precipitates occurred in
the thermomechanically affected zone (TMAZ) and the HAZ. Complete dissolution of coarse γ′ precipitates was
evident in the stir zone (SZ).
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18 Heat-Affected-Zone Degradation of Mechanical Properties
Figure 18.29 Linear-friction-welded Waspaloy rectangular rods (13 mm × 11 mm in cross-section): (a) workpiece after
welding; (b) hardness profiles in as-welded (As-LFWed) and postweld heat-treated (PWHTed) conditions; (c) primary (P),
secondary (S) and tertiary (T) precipitates of γ′ in base metal; (d) dissolution of precipitates at weld interface. Source:
Chamanfar et al. [41]. © Elsevier.
Figure 18.29 shows of the dissolution of γ′ precipitates at the weld interface in linear friction welding of
rectangular Waspaloy rods [41]. An in-plane actuator oscillated the lower part in the horizontal direction
and a forge actuator applied a downward force on the stationary upper part. The forging pressure was 360 MPa
and the axial shortening was 4.6 mm. The composition of the Waspaloy was Ni-18.08Cr-12.87Co-1.00Fe4.12Mo-3.35Ti-1.07Al-0.07B-0.01 Zr-0.04C in wt%. The alloy was welded in the as-received condition of solution heat treating and double aging. The base metal shows primary (P) secondary (S) and tertiary (T) γ′
precipitates. The as-welded hardness profile (triangles) shows significantly softening at the weld interface.
This is caused by dissolution of γ′ precipitates at the weld interface. The hardness profile (circles) shows
recovery of hardness after postweld heat treating. For postweld heat treating, the welded workpiece was furnace heated to 1000 °C, soaked at this temperature for 1 hour, and then air cooled to room temperature. This
was followed by a double aging procedure consisting of furnace heating to 850 °C, holding for 4 hours followed by air cooling to room temperature, then reheating to 760 °C, holding for 16 hours and then air cooling
to room temperature.
Damodaram et al. [42] showed of the dissolution of γ″ precipitates at the weld interface in rotary friction
welding of IN 718 rods of 13 mm diameter, with the composition of Ni-18.2Cr-5.1Nb-3.28-Mo-1.06Ti-0.56Al0.33V-0.09Mn. The rods were heat treated in two different conditions before welding. Welding was conducted
18.4 ­Dissolution of Precipitates in Ni-Base Alloy
Figure 18.30 Hardness profiles in IN 718 welds in as-welded condition: (a) laser welds; (b) GTA welds. S: solutionized; A:
aged after solutionization; L: laser welded; T: GTA welded. Broken line indicates fusion line. Source: Hirose et al. [43] (36). ©
University of Sheffield.
Figure 18.31 Hardness profiles in IN 718 welds after postweld heat treating: (a) laser welds; (b) GTA welds. S: solutionized;
A: aged after solutionization; L: laser welded; T: GTA welded. Broken line indicates fusion line. Source: Hirose et al. [43] (36).
© University of Sheffield.
in the solution treated (ST) condition, which was done at 995 °C for 1 hour. Welding was also carried out in the
solution treated and aged (STA) condition, in which the solutionized alloy was aged at 720 °C for 8 hours followed by furnace cooling to 620 °C and then aging at 620 °C for 8 hours followed by air cooling to room temperature. Welding in the ST condition did not cause softening near the weld interface. In fact, the hardness
increased slightly near the weld interface due to grain refining caused by dynamic recrystallization. Welding in
the STA condition, however, resulted in significant softening near the interface. This was caused by dissolution
of γ″ precipitates near the weld interface. The hardness could be recovered by direct aging (DA) after welding
regardless of whether welding was conducted in the ST or STA condition. The aging condition was 720 °C for 8
hours, followed by furnace cooling to 620 °C and then aging at 620 °C for 8 hours before air cooling to room
temperature.
Hirose et al. [43] studied the effect of the welding process on softening in the HAZ of IN 718 by comparing laser-beam welding with GTAW. As shown in Figure 18.30, when IN 718 is solutionized (S) to dissolve
γ″ precipitates before welding, neither laser-beam welding (SL) nor GTAW (ST) shows softening in the
HAZ. However, when IN 718 is aged before welding, GTAW (AT) shows a much wider HAZ of softening
than laser-beam welding (AL). The higher heat input in GTAW can be expected to cause more dissolution
of γ″ precipitates. Figure 18.31 shows the hardness can be recovered by post-weld aging, either direct aging
or solution heat treating followed by aging [43].
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18 Heat-Affected-Zone Degradation of Mechanical Properties
18.5 ­Martensite Tempering in Dual-Phase Steels
Figure 18.32 shows the microhardness profile across a weld of dual phase steel DP600 (tensile strength
600 MPa) made by laser-beam welding and SEM images of selected areas [44]. The rapid cooling during welding results in the formation of martensite in the fusion zone and a considerable increase in the hardness.
However, a soft zone forms in the HAZ due to the tempering of the preexisting martensite in the DP600 steel.
This soft zone corresponds to the subcritical region of the Fe-C phase diagram, as shown in Figure 18.33. This
region is just below the A1 temperature of the Fe-C phase diagram. Thus, the location of the soft zone in a dual
phase steel is equivalent to the area just outside the HAZ of an ordinary carbon steel. In other words, for dual
phase steels the HAZ is widened to include subcritical region of the Fe-C phase diagram. By both SEM microstructure examination and nano-hardness testing, Hernandez et al. [45] showed martensite is weakened when
tempered.
Figure 18.34 shows the microhardness profile across a weld of dual phase steel DP980 made by laser-beam
welding and SEM images of selected areas [44]. As compared to DP600 (Figure 18.32), the volume fraction of
martensite is significantly higher, and this contributes to its higher strength (980 MPa). Again, fast cooling during
laser-beam welding causes much martensite to form in the fusion zone and results in a higher hardness level in
the fusion zone. However, the extent of softening and the size of the soft zone are significantly larger in DP980
than in DP600. Thus, in welding dual phase steels of a higher strength level, more significant HAZ weakening due
to martensite tempering can be expected.
Figure 18.32 Laser weld of dual phase steel DP600: (a) microhardness profile; (b) SEM image of base metal; (c) SEM image
of soft zone; (d) SEM image of fusion zone. Source: Farabi, Chen, and Zhou [44]. © Springer.
18.5 ­Martensite Tempering in Dual-Phase Steel
Figure 18.33 Subcritical region of dual phase steels where martensite can be tempered: (a) peak temperature profile; (b)
Fe-C phase diagram.
Figure 18.34 Laser weld of dual phase steel DP980: (a) microhardness profile; (b) SEM image of base metal; (c) SEM image
of soft zone; (d) SEM image of fusion zone. Source: Farabi, Chen, and Zhou [44]. © Springer.
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18 Heat-Affected-Zone Degradation of Mechanical Properties
Examples
Example 18.1 2219 Al-T6, 5 mm thick, was bead-on-plate welded by friction stir welding at 800 rpm tool rotation speed and 200 mm/min travel speed. The shoulder diameter was 14 mm and the pin length 4.5 mm. Overaging
was reported. (a) Which precipitate phase is likely to dominate in the base metal? (b) Why did overaging occur?
(c) Where was overaging most significant, in the stir zone, thermomechanically affected zone (TMAZ) or HAZ?
(d) Which phase existed in the region of most significant overaging?
Answer
(a) The θ′ phase. The TEM image show fine metastable θ′ in the base metal. (b) Overaging occurred because the
2219 Al was welded in the T6 condition, i.e. with the optimum strength (hardness). (c) In the stir zone and the
TMAZ, where the highest temperature was reached during FSW. This was confirmed by the microhardness
distribution across the weld. (d) The θ phase. TEM images showed coarser equilibrium θ in both the stir zone
and the TMAZ. Check Chen et al. [19, 20] for more details.
Example 18.2 A work-hardened Al-Mg alloy 5083-H321 was friction-stir welded. The hardness profile across
the weld showed, on the HV 2.5 scale, the hardness dropped from about 90 in the base metal to about 75 near the
weld centerline. An annealed Al-Mg alloy 5083-O welded under the similar conditions, on the other hand, showed
the hardness increased from about 72 in the base metal to about 77 near the weld centerline. (a) Explain the hardness drop in 5083-H321. (b) Explain the hardness increase in 5083-O.
Answer
(a) The hardness drop in 5083-H321, which is in the work-hardened state, is caused by softening due to
recrystallization (dynamic). (b) The hardness increase in 5083-O, which in the lowest hardness state, is
due to the finer grains in the stir zone produced by recrystallization (dynamic). According to the HallPetch equation, Eq. (18.1), the strength and hence hardness are higher with finer grains. For more details,
see Threadgill [46].
­References
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2 Yu, P., and Kou, S. Research in progress. University of Wisconsin, Madison, WI.
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molybdenum alloy. Materials Science and Engineering 59 (2): 257–273.
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and Titanium Alloys at Lawrence Livermore National Laboratory using a Fiber Delivered 2.2 kW Diode Pumped
CW Nd: YAG Laser. 2006, Lawrence Livermore National Laboratory (LLNL), Livermore, CA.
7 Tabernig, B. and Reheis, N. (2010). Joining of molybdenum and its application. International Journal of Refractory
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8 Zhang, P., Li, S.X., and Zhang, Z.F. (2011). General relationship between strength and hardness. Materials Science
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9 (1972). Metals Handbook, 8e, vol. 7, 135. Metals Park, OH: American Society for Metals.
References
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11 Cook, L.A., Channon, S.L., and Hard, A.R. (1955). Properties of welds in Al-Mg-Mn alloys 5083 and 5086. Welding
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14 White, S.S., Manchester, R.E., Moffatt, W.G., and Adams, C.M. (1960). Plastic properties of aluminum-magnesium
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Materials Science and Technology 19 (6): 785–790.
16 Kolarikova, M., Kolarik, L., and Vondrous, P. (2012). Welding of thin molybdenum sheets by EBW and GTAW. In:
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17 Dumolt, S.D., Laughlin, D.E., and Williams, J.C. (1982). The effect of welding on the microstructure of the age
hardening aluminum alloy 2219. In: Proceedings of the First International Aluminum Welding Conference. New
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18 Chen, Y.C., Feng, J.C., and Liu, H.J. (2009). Precipitate evolution in friction stir welding of 2219-T6 aluminum
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19 Chen, Y., Liu, H., and Feng, J. (2006). Friction stir welding characteristics of different heat-treated-state 2219
aluminum alloy plates. Materials Science and Engineering A 420 (1): 21–25.
20 Liu, H.J., Zhang, H.J., and Yu, L. (2011). Effect of welding speed on microstructures and mechanical properties of
underwater friction stir welded 2219 aluminum alloy. Materials & Design 32 (3): 1548–1553.
21 Dumolt, S.D., Laughlin, D.E., and Williams, J.C. (1981). The effect of welding on the microstructure of the age
hardening aluminum alloy 2219. In: Proceedings of the First International Aluminum Welding Conference, 115. New
York, Cleveland, OH: Welding Research Council.
22 Lin, Y.T., Wang, D.P., Wang, M.C. et al. (2016). Effect of different pre-and post-weld heat treatments on
microstructures and mechanical properties of variable polarity TIG welded AA2219 joints. Science and Technology
of Welding and Joining 21 (3): 234–241.
23 Kou, S., and Le Y. (1982). Unpublished research. Carnegie Mellon University, Pittsburgh.
24 Soysal, T., Yu, P., and Kou, S. (2017). Research in Progress. Madison, WI: University of Wisconsin.
25 Liu, H.J., Zhang, H.J., Huang, Y.X., and Lei, Y.U. (2010). Mechanical properties of underwater friction stir welded
2219 aluminum alloy. Transactions of Nonferrous Metals Society of China 20 (8): 1387–1391.
26 Lei, X., Deng, Y., Yin, Z., and Xu, G. (2014). Tungsten inert gas and friction stir welding characteristics of
4-mm-thick 2219-T87 plates at room temperature and− 196° C. Journal of Materials Engineering and Performance
23 (6): 2149–2158.
27 Martukanitz, R.P. and Howell, P.R. (1996). Trends in Welding Research, 553. Materials Park, OH: ASM
International.
28 Kou, S. and Le, Y. (1985). Improving weld quality by low-frequency arc oscillation. Welding Journal 64 (3): 51s–55s.
29 (1964). Metals Handbook, 8e, vol. 2, 276. Metals Park, OH: American Society for Metals.
30 Kaufman, J.G. (2000). Introduction to Aluminum Alloys and Tempers, 39–76. Materials Park, OH: ASM
International.
31 Metzger, G.E. (1967). Some mechanical properties of welds in 6061 aluminum alloy sheet. Welding Journal 46 (10):
457s–469s.
32 Sato, Y.S., Kokawa, H., Enomoto, M. et al. (1999). Precipitation sequence in friction stir weld of 6063 aluminum
during aging. Metallurgical and Materials Transactions A 30 (12): 3125–3130.
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33 Burch, W.L. (1958). The effect of welding speed on strength of 6061-T4 aluminum joints. Welding Journal 37:
361s.
34 Hirose, A., Kurosawa, N., Kobayashi, K.F. et al. (1999). Quantitative evaluation of softened regions in weld
heat-affected zones of 6061-T6 aluminum alloy. Metallurgical and Materials Transactions A 30 (8): 2115–2120.
35 Zhou, Z. (1981). Principles and Technology of the Fusion Welding of Metals, vol. 1, 157. Beijing, China: Mechanical
Engineering Publishing Co.
36 Mizuno, M., Takada, T., and Katoh, S. (1967). Weldability of Al 4.5%Zn 1.2%Mg alloy. Journal of the Japan Welding
Society 36 (8): 854–861.
37 Kalemba, I., Hamilton, C., and Dymek, S. (2014). Natural aging in friction stir welded 7136-T76 aluminum alloy.
Materials and Design 60: 295–301.
38 Zhang, K., Chen, J.Q., Ma, P.Z., and Zhang, X.H. (2018). Effect of welding thermal cycle on microstructural
evolution of Al–Zn–Mg–Cu alloy. Materials Science and Engineering A 717: 85–94.
39 Owczarski, W.A. and Sullivan, C.P. (1964). A microstructure study in a welded superalloy (welding effect on
nickel-base superalloy, examining microstructural changes). Welding Journal 43 (9): 393s–399s.
40 Jun, H.J., Ayer, R., Neeraj, T., and Steel, R. (2007. Trans Tech Publications). Friction stir welding of precipitation
hardened Ni based alloys. Materials Science Forum 539–543: 3763–3768.
41 Chamanfar, A., Jahazi, M., Gholipour, J. et al. (2012). Maximizing the integrity of linear friction welded Waspaloy.
Materials Science and Engineering A 555: 117–130.
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alloy 718. Materials Science and Engineering A 560: 781–786.
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Engineering Materials, University of Sheffield.
44 Farabi, N., Chen, D.L., and Zhou, Y. (2012). Tensile properties and work hardening behavior of laser-welded
dual-phase steel joints. Journal of Materials Engineering and Performance 21 (2): 222–230.
45 Hernandez, V.B., Panda, S.K., Kuntz, M.L., and Zhou, Y. (2010). Nanoindentation and microstructure analysis of
resistance spot welded dual phase steel. Materials Letters 64 (2): 207–210.
46 Threadgill, P. (1977). Friction stir welds in aluminum alloys – preliminary microstructural assessment. TWI
Bulletin.
Further Reading
Lu, Y., Peer, A., Abke, T. et al. (2018). Subcritical heat affected zone softening in hot‐stamped boron steel during
resistance spot welding. Materials & Design 155: 170–184.
Problems
18.1
A 12.7 mm thick plate of 6061-T6 aluminum (TL = 652 °C) was GTA welded with DC electrode negative.
The welding parameters were I = 222 A, E = 10.4 V, and V = 5.1 mm/s. Microhardness measurements after
welding indicated that softening due to overaging starts about 5.3 mm from the fusion line and gradually
increases as the fusion line is approached. Thermal measurements during welding revealed a peak temperature of about 300 °C at the position where softening started. Calorimetric measurements showed that
the arc efficiency was around 80%. How does the width of the HAZ compare with that predicted from
Adams’ equation (Chapter 2)?
Problems
18.2
Al-Li-Cu alloy 2095 was welded by LBW, GTAW, and GMAW and the HAZ hardness profiles of the resultant welds were measured. Rank the welds in the order of increasing hardness in the HAZ.
18.3
Al-Li-Cu alloy 2090 was welded with a matching filler metal. Joint efficiencies up to 65% of base metal
strength were obtained in the as-welded condition. After postweld solution heat treatment and artificial
aging, joint efficiencies up to 98% were obtained. Explain why.
18.4
7075-T6 Al (base metal hardness about 165 on the HV 5 scale) was treated in the following three different
ways after welding: (i) natural aging for three months at room temperature (pwna), (ii) artificial aging
(pwaa), and (iii) full post weld heat treatment of solutionizing, quenching and then artificial aging (pwsh).
(a) In Case (i) the HAZ hardness was about 140 near the fusion line (up to 8 mm from the fusion line),
decreased gradually to a minimum of 120 at about 12 mm from the fusion line, and increased gradually to
165 at about 33 mm from the fusion line. Explain the hardness distribution. (b) In Case (ii) the minimum
hardness was about 100 at about 12 mm from the fusion line. Explain why the minimum hardness was
lower in Case (ii) than in Case (i). (c) The hardness was uniform at about 165 across the HAZ and the base
metal. Explain why.
503