Mg Alloy Dev Guided by Thermo Calculations
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Transcript of Mg Alloy Dev Guided by Thermo Calculations
MAGNESIUM ALLOY DEVELOPMENT GUIDED BY THERMODYNAMIC CALCULATIONS
Joach~m Grobner. Dmytro Kevorkov, Rainer Schmid-Fetzer
Technical University of Clausthal. Institute of Metallurgy, Robert-Koch-Str. 42
D-38678 Clausthal-Zellerfeld. Germany
Abstract
In tradit~onal alloy development. expenmental investigations mith
many different alloy compositions are performed. The selection
criter~a for multicomponent alloying elements and the~r
cornpos~tions become diffuse in a traditional approach.
Computational thermochemistry can provide a clear g ~ ~ i d e l ~ n e for
such selections and helps to avoid large scale experiments with
less promising alloys. Thus. ~t is a powerful tool to c ~ ~ t down on
cost and time during development of Mg-alloys. As an example.
me report on recent developments of nem, creep resistant alloys
that show about 100 times less creep than the best con~mercial
alloys. Also outlined are the methods we use in our long-term
project of construction of the necessary therrnodynam~c database
for several prom~sing alloying elements like Al. Li. SI. Mn. Ca.
Sc. Y. Zr and Rare Earths. using the Calphad method cornb~ned
~ ~ t h key experiments.
Introduction
There is an urgent need for the development of new or improved
magnesium alloys if \4e want to fully exploit the potent~al of this
t'asc~nating lightue~ght mater~al that also offers excellent
cxstah~llty, ma china hi lit^ and b~o-compat~bility. Experiments on a
technological scale f o r preparation and testing o f new a l l q s are
\er> eupensl\e and time consuming. In vie^ of the huge number
oi' posvhle allo) components. compositions and processing
parameters. one \\auld like to hale at least an "educated guess" in
which direction to go. In this report we want to s h o ~ that
thermodynamic calculations can provide much more than that.
Computational thermochemistry is a modem tool that supplies
quantitative data to guide the development of alloys. It enables the
calculation of rnulticomponent phase diagrams and the tracking of
individual alloys during heat treatment or solidification by
calculation of phase distributions and phase compositions. These
are the basic data to understand and control the behavior of any
novel or modified Mg-alloy. Large-scale experiments for new
multicomponent alloys can then be focused on most promising
alloys identified in that approach. Long-term experiments with
less promising alloys can be avoided. Thus, it is a pom'erful tool to
cut domn on cost and time during development of magnesium
alloys.
Database Develournent
In the core of this method a h ~ g h quality thermodynamic database
for Mg-alloys is needed. Development of a reliable
thermodynamic database for multicomponent alloys requlres a
combination of experiments and computational thermochenlistrq
with data Prom alloy application. Since numerous binary and
ternary subs)stenis have to be treated and validated \v~th key
experiments before multicomponent alloys can be calculated
reliably. t h ~ s development becomes a long-term project. In our
group at the TLI Clausthal a thermodynamic database for seberal
alloying elements like Al. Li. Si. Mn, Ca. Sc. Y. Zr and Rare Earth
elements (Fig. I ). has been under construction for more than fi \e
bears and is still ongoing.
Magnwum Technologj 2001 Edited by J Hr?n
TMS (The Mnerais. Metals & Materials Society). 2001
I Creep resistance - Density reduction I ' \
Fig. I: Currently selected alloying elements in the Mg-database.
To create the thermodynamc phase descriptions the Calphad
method is used. The principle of this method IS shoan in Fig. 2
Thermodynamic Modelling
Fig 2 Schemat~c approach of database development and the Imks to
appl~catlons
In a multicomponrnt, niult~phase system each phase is described in a
suitable model for the G h b s energy. The parameters are optimi~ed
by avmlable expenmental data integrating both thernwdynamic and
phase equil~bnurn data. Tn_ls data set is used to calculate stable and
neta astable equi l~hr~a for process simulat~on or direct applmt~ons.
The calculat~ons are ~ m p r o ~ e d cont~nuously by new espcnrnental
data coming from apphcations nnd are checked hy key evperlments
to improve their r e l ~ a b ~ l ~ t y . An example for process simulat~on 1s
the coupl~ng u ~ t h 3D-sol1d1ficat1on n~odel~ncp. Example\ for d~rect
applicat~ons are given in the followng chapters.
In Mg-system5 the exper~mental l ~ t e r ~ t u r e data \+ere \ery q w s e
In both. phase d~agram data and the rmod~nam~c proprrt1e5
Therefore, seheral exper~mental methods are b a n g applied in our
group to produce a sufficient exper~mental database (Fig. 3) .
Experimental Methods
Sample Preparation Levltatlon Meltlng Arc Meltlng Electron Beam Meltlng Reaction Slnterlng
Sample Analysis ThermalAnalys~s (DTA,STA) X-ray Dlffractornetry ( X R D , 25-1500" ) Electron Mlcroscopy (SEMIEDXIWDX) Metallography, Opt~cal Mlcroscopy
Fig. 3: Experimental methods.
Our ao rk focuses on ternary and quaternary systems for
~mprovement of creep reslstance [ I ] , thermal s t a b ~ l ~ t ) zpra?
forrmng (Mg-Mn-Sc-RE. Mg-AI-Ca-RE) and density reduct~on
(M~-LI-AI-SI. Mg-AI-Ca-Si) [2. 3, 4. 5 . 61 The thermodynamc
and technical application of the temarj Al-Mg-Sc 15 alreddb
discused In l~terature [7, 81
Here the quaternary systems Mg-Mn-Sc-Gd and Mg-Mn-Sc-Y are
shown as example for the selection of new creep resistant alloys
uslng computational thermochem~stry.
First generation of creep resistant ternary hIg-Mn-Sc alloys
In~est~gat ions started w t h binary Mg-Sc alloys. Scandium \\as
chosen for precipitation hardenmg because of its large solub~lity
In (Mg) and the retrograde solubility at louer temperatures. The
b~nary phase d~agram had to be re-investigated (91. it is s h o ~ n in
Fig. 3. Ho\+ever, hinary MgSc precipitates form Lery s lo~. l !
dur~ng ageing and improve the mechanical properties on14
slightly because of their incoherent interface. Therefore. Mn uas
added as second alloy~ng element. The precipitatmn of Mn&
\+as predicted by thermodynamic calculat~ons Mn&
precip~tatmns form coherently and uere found very useful for
~rnproi ing creep reslstance and hardness. Neu, MgSc l5Mn l or
MgSc6Mnl alloys shou about 100 tlmes better creep resistance
than the best commercial IY E43 alloy at 350°C and 30 hIPa [ I ]
In spite of the good properties of t h ~ s first generation of h,lg-hln-
S c alloys. the high cost of S c add~t lon ( 6 wt.'? S c o r more)
~ n ~ t ~ a t e d a search for a second generation by i n v e s t ~ g a t ~ n g
quaternary systems.
At this point the questlon arose. "ho14 to ldent~t:, pronilslng a l l q
cand~dates from all these calculated diagrams." What phase
diagram features are related lo what alloy processing steps .' Whal
is needed is a list of beneficial phahe diagram feature\. d r r ~ ~ e d
from the relevant alloy processing steps. The most Important
polnts are g l w n In Table I
0 l I I I I I I
0 2 0 4 0 6 0 80 1 0 0 M g at . O 6 S c
s c
FIE 1 Thc b ~ n a ) Mg-SL \)\tern atter [9]
A d d ~ t ~ o n a l a l loyng elements Gd. Y and Zr were considered for
t h ~ \ purpose to achieve a sufficient quantity of suitable
preclpltatcs to Improve rncchan~cnl propertics u u n g n mlnlmum
(>I' eupensi\e alloy element add~t ion . These elerncnt comb~nat ions
hlg-MII-(Sc. Gd. Y. Z r ) f o ~ m a i a n e t y o l quatcrnar) systems and
thin tho\c there is a huge amount o f p o s s i h i l ~ t ~ e ~ to select allog
compo\itions. Therefore phase d~ayrai i i and other cillculatlons
\\?I-e performed to ~ d e n t ~ f ) p ro in i~ ing c a n d d a t t s
Ulo\ selection in the 51~-\In-Sc-(;d ,\stem
5 0 G d 0.2 0.4 0.6 5 0 G d
1 O M n W t . % S C 1 O V n 9 4 0 M g 9 3 0 M g 0 0 S c 1 osc
F I ~ . 5: Phase d ~ a g r a m section n ~ r h constant I n t i; Mn. 5 U I
Table 1: Beneficial phase equilibrium features and their
relevance for (5Ig)-alloy processing
I Phaw diagram feature I Kele\ance for alloy
Larpe t.nougti c Alp) 51nglc- Enable\ Iiomogeni/ation plia\e field ailnealing
For an alloy ~vith I wt.Q Mn. 5 nt.% Gd and 0 8 ut. '? Sc
(indicated by arrow In Fig. 5 ) equ~ l ib r~um phase amoilnts during
sol~d~fication and heat treatment are piken In Fig. 6. At the
I ~ q u ~ d u s temperature 650°C. prlrnary (Mg) is formed and
consumes the melt totally up to the sohdus polnt at 6 1 9 T At
590°C the first prcclpltate Mn2Sc starts forrmnp, which can he
seen In the enlarged Insert in Fig. 7. Large amounts of the second
preclpltate Gdhlg, start forming at 325'C. This alloy fulfills all
features illustrated In Table 1 and was classified as ver)
promising for further alloy development. Similar features are
ohserved for an alloy with even less scand~um (0 3 wt.G Sc) .
In h c t , first results for this alloy shows a creep resistance similar
to the prevlous ternary high-scandium Mg-Mn-Sc alloys, about
100 times better than best commercial WE33 alloy at 350cC and
30 MPa as detailed later.
temperatures. In the whole ranye ot primary cr) \ ta l l~/a t~oi i oi
(My) the secondary phase is Mnl lY. The des~red LIn,Sc
preclpltate f'orms only at temperatures he lo^ 500'C. wh~cti
decrease ~ ~ t h decreas~ng Sc c o n m t .
0 1 I I I
0 0 2 0 4 0 6 0 8 1 0
P h a s e amoun t [mo l l
Fig. 6 : Phase amounts of MgMnIGdiSc0.8 alloy
Alloy selection in the hIg-Mn-Sc-Y system
In the Mg-Mn-Sc-Y system heveral ~er t ica l \ectlon\ In the I-anyer
01'0-1.5 L I ~ . ' / ; Mn. 0-10 wt.'; Sc and 0-10 ut . ' i Y \\ere stud~ed
F I ~ 7 ihous a T-x sectlon of a calculated cjuaternar> phase
e q i ~ ~ l ~ h r i a ~11th constant I u t C; M n . 5 \\t.':; 1' from 0- I \ i t . ' ; Sc.
A large one-pha\e ticld o f ( M g ) and se\e~-a1 d~ffcrcnt ro l~d p h a m
\table at Iouer ternpcratures can be seen, like in the con-espond~ng
d~agrani f'or (id. Diftercnces cornpal-cd to the Gd \)stern can he
ohwr\cd concernln? the htablc so l~d phases and the forniat~on
Fig. 7: Phase d~ayrani section ulth constant I ~ t . 7 4 hln. 5 L I ~ ' i Y
from 0- 1 w t . 5 Sc.
200
MgSc 0
liquid
Mg24Y5
0 06 0.08 0 1
0 0 2 0 4 0 6 0 8 10
Phase amount [moll
Fig 8 . Phase anioilnts of LIgLln I1'5Sc0.8 allo!
Pol- an alloy \ \ ~ t h I \\ t ' i ivln. 5 i \ t r'i 'r' and 0.8 M t ' i Sc
I ~ n d ~ c a i e d b) arrou In Fi?. 7 ) phaw amounts durlng soiid~licat~on
ai. gl\cn In F I ~ . 8. Xr [he 11qu1du\ temperature 644°C. prirnar!
I 1 1 ~ ) 1s formed and consumes the ~ ~ i e l t totally up to the sol~dus
po~nt at 613-C. The Mn12\r' phase would onl! be stable In the
\()I111 \tat? In a high temperature range from 605 to 322°C and
n i~ ! not form at all du r~ng fast cool~ng. The estahllshcd Mn2Sc
and the ncu hlg24'r'5, eien In \cry large amount, could form by
agtlng in a favorable temperature range. T h ~ s makes
LlyhlnI Y5ScO.S also a promis~ng alloy.
In 01-der k-r checA if the amount of MnJc could he raised h j
Increasing the manganese content to 1.5 u t ', u e can examine
F I ~ 9 J'or the alloy MpMnl.SY.5Sc0.8. Thls alloy I S d~squalilied
LIncc Mn2Sc forms a j the prlrnar) phase from the I ~ q u ~ d . clearly
wen In thc Inset of F I ~ . 9. (Mg) fornis only secondar!. and e \en
Lln,:\r' limns as a terliary phase du r~ng solld~fication. Such a
nilcrwtructure cannot be "repalred" h! anneal~ng
cn 0 0 0 2 0 0 4 0 0 6 0 0 8 0 . 1
0 1 I I I
0 0 2 0 4 0 6 0 8 1 0
P h a s e a m o u n t [ m o l l
i l lo, \election in the \lg-hln-'l -%I. \\stern
u~bstlt~ltlon can he srudled In Flg. 10 for scandli~ni frt'c ;lllo)<
iuth 1 ~ 4 t . q ' Mn. 4.5 i 4 t . q Y and 0-1 u t . ' i Zr. T h ~ s phase
d~agrarn sectlon sho\<s a ver) steeply rlslng I ~ q u ~ d u s lint. l000'C
are reached for less than 0 .1 wt.% Zr. Moreover. and actually the
reason for the steep liqu~dus line. a huge primary crystali~zat~on
tield. L + Mn2Zr. stretches over the entire composition range In
Fig. 10 Only for extremely small Zr-add~tions. not d~sccrnible In
Fig. 10. is a prlmary (hlg) sol~diiicat~on expected The reason for
thls destructi\e phase dlagram f e a t ~ ~ r e I S the extremely high
thermodynamic stabihty of .MnzZr in cornparlaon to the other
phases. Since yttrlurn does not play a significant role In that part.
the only \\,a)' to dimrn~sh the L+Mn2Zr prlmar} ficld ~~l.ould he ;I
drastic reduct~on of the manganese content. But this \\auld also
drastically reduce the amount of beneficial Mn-conta~n~ng
preclpltates. As a result. the entire quaternary Mg-Mn-Y-Zr alloy
\) ztem 1s disqualified.
F I ~ . 10: Phase d l a ~ r a m sectlon wtli con,tant I \ \ t c; Xln. 5 \ \ t
1' lrom 0- I \ 4 t 3 Zr.
- 2 -
111t goal ofdrclaticall)~ rrduclng the Sc-contc~il c m he ac~cc~rnpl~hcd
More cletails of alloy selection In the q~iate~nary s) \ kms 112-kln-
.Sc. Gd. Y. Zrj \ \11l he publishecl \eon [ 101
L t ( M g ) + M n ' Z r
Comparison to commercial creep resistant Mg alloys Alloy preparation and creep resistance measurements
In order to assess the results on the creep resistance of the selected
new alloys we have to choose a commercial alloy as a benchmark.
The history of commercially developed creep resistant alloys is
briefly summarized in Fig. 11.
Alloy Composition (mass%) Year Zn A g Y R E
- most of these alloys conlain as gram refmer about 0 7% Zr - R E , the inf luence decreases In the order Nd (Pr) > C e > La
Fig. 1 1: Commercial creep resistant Mg-alloys
Alloys like AE42 or AS21, known to be more creep resistant as
standard AZ or AM alloys, are not included in the listing of Fig.
I I since they rank below the ZE alloys. The current end point of
commercial alloys is given by the WE series containing Y and
Rare Earths. In fact, the maximum stress that can be tolerated at
200°C for 100 h and 0 . 2 6 elongation is highest for the WE series
as seen in Fig. 12. The alloy WE43 with T6 heat treatment was
chosen as the benchmark alloy.
185 181 Lord 100 h, 200'C
WE54 18 WE43 16 OE22 18 ZE41 T5 A281 T4
Alloy
Fig. 12: Creep properties of commercial magnesium alloys.
The most promising alloy compositions identified by the
thermochemical calculations were prepared by squeeze casting by
our project partners within the Thrust Research Project SFB 390.
"Magnesium Technology", at the TU Clausthal. Yield strength
and creep rates of these alloys were measured in the as cast
condition and also after heat treatment [I I ] . It is clearly seen from
Fig.13 that the secondary creep rate of our first generation (high-
scandium) alloys is better by a factor of 100 compared to the
WE43 benchmark alloy.
T=350°C, -30 MPa 1 .SlI
time [s] Fig. 13: Creep curves of first generation (high scandium) alloys
compared to the WE43 benchmark [I 11.
The first generation of alloys showed a strong annealing response
due to the formation of the Mn,Sc precipitates. The existence of
the MnzSc precipitates was confirmed by X-ray diffraction. SEM
and TEM investigations and energy dispersive X-ray
microanalysis. A micrograph showing finely dispersed Mn&
precipitates in a MgSc6Mnl alloy after T5 treatment is given in
Fig. 14.
Conclusions
Fig. 14: Micrograph of MgSc6Mnl T5 alloy.
The second generation (low-scandium) alloys were selected from
the most promising candidates identified by the thermodynamic
calculations given in the previous sections.
The creep curves for these low-scandium quaternary alloys are
shown in Fig. 15. Again, an almost 100 times lower creep rate is
achieved at 35OoC and 30 MPa compared to the best commercial
alloy WE43. The best alloy MgGdSMnlSc0.3 with the smallest
elongation after 1 5 . 1 0 ~ ~ contains only 0.3 w t . 8 of expensive Sc.
T=35OoC, -30 M P a
'O time [ lo4 s] l5
F I ~ 15: Creep curves of second generation low-scandium alloys
compared to the WE43 benchmark.
This I S substantial progress compared to the first generation of
alloys with 6 or 15 wt.% Sc. Further work on the selected alloys
is in progress.
Focused magnesium alloy development is now possible
using the powerful tool of thermodynamic calculations.
Alloy compositions with promising possibilities of alloy
microstructure design can be selected by means of
thermodynamically calculated phase diagrams, phase
amount charts and solidification curves. Most importantly,
element combinations and compositions with unwanted
properties can be recognized before starting large-style
experiments, thus reducing the experimental effort to a
reasonable volume. The next step, the experimental study
of mechanical properties of identified promising alloys has
shown excellent results both in the first generation
(ternary) and second generation (quaternary) of new creep
resistant alloys.. Obviously, these experiments cannot be
replaced by thermodynamic calculations. However,
considering the huge number of less promising alloy
combinations that could have been selected from
multicomponent systems, the focused alloy development
following this approach avoids a waste of time and effort.
Acknowledgement
This work is supported in the thrust Research Project SFB 390:
Magnesium Technologya by the German Research Council
(DFG).
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