application of electron back scatter diffraction
TRANSCRIPT
Electron Diffraction Subcommittee 22 March 2007 Ray Goehner
Call to Order R. Goehner provided comic relief in the form of a brief Dilbert cartoon. Appointment of Minutes Secretary T. Kahmer Board of Directors’ Liaison Report Motions presented to the board in March 2006, and their responses. - Not read. Confidential information on record at ICDD. Business R. Goehner gave a presentation of how electron diffraction is used for phase identification. Presentation follows the minutes. Backscattered Electron Kikuchi Patterns BEKP Electron Backscatter Patterns EBSP Backscattered Kikuchi Diffraction BKD Wide Angle Kikuchi Patterns WAKP T. Fawcett will contact the editor of International Materials Reviews and ask for permission to put a copy of the article, “Application of electron backscatter diffraction to the study of phase transformations”, 2007 vol. 52 no. 2, pp 65-128, on the ICDD web site. Article follows the minutes. There was some discussion on primary and alternate codes, particularly in the case of silicon dioxide because of its many polymorphs. J. Friel volunteered to be our spokesperson for the community. He will try to get a focal group of ~six people together for a conference call to discuss how to get this project started. There is a meeting in aluminum industry in May. J. Friel will most likely get to speak with Scott Sitzman of HKL and Stuart Wright from EDAX. HKL has a lot of information on their website that shows techniques and it explains the phase map, i.e., different colors with different phases. They offer three databases with their product. T. Fawcett and J. Faber talked to Dawn Janney (a new ICDD member from Idaho National Laboratory). She bought a PDF-4+ and was excited about it. She gave input in how to enhance the product for electron diffraction. Adjournment
EBSD in the SEM for the Identification of Unknown Crystalline Phases
DXC 2004 WorkshopDXC 2004 WorkshopPrinciples & Use of Principles & Use of Microdiffraction &Microdiffraction &MicrofluorescenceMicrofluorescence
Ray Goehner & Joe MichaelSandia National LaboratoriesAlbuquerque, NM 87185-0886
Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company,for the United States Department of Energy under contract DE-AC04-94AL85000.
EBSD Pattern of AlB2
Pattern Features:
Parallel lines are Kikuchi line pairs
Spacing between pairs is twice the Bragg angle and inversely related to the d-spacing
Places where lines intersect is called zone axis
Angles between zone axes are indicative of crystal structure
Angles made by Kikuchi line pairs about a zone axis are the interplanar angles of the crystal
AlB2 (20 kV)
History of EBSP
Backscattered Electron Kikuchi Patterns
BEKP
Electron Backscatter Patterns
EBSP
Backscattered Kikuchi Diffraction
BKD
Wide Angle Kikuchi Patterns
WAKP
First observed in 1954 - before SEM invented
1970’s Venables et al. were first to observe EBSP in SEM
1980’s Dingley et al. began using these patterns for orientation studies
1990’s Wright and Adams developed automatic system for texture (OIM)
Michael & Goehner. develop technique for phase identification
Cameras for EBSD – Typical Arrangement
Sample surface Phosphor screen
Pole piece
Not much room in the sample chamber!
Installation on a Dual Beam FIB
Most SEM sample chambers present challenges to installation of EBSD cameras.
But, it can usually be done with a little bit of compromise.
Effect of Sample Tilt on BS Electron Yield
# B
acks
catte
red
elec
tron
s
0 10 20 30Energy (kV)
Tilted sample has higher BS electron yield
Sample tilt results in sharp peak in BS electron energy distribution. Better defined energy of BS electrons results in sharper Kikuchi lines.
Alumina (Al2O3) On Thin Substrate
Ag on MgO – Example of EBSD Resolution
Orientation map of Ag particle normal direction
SEM images of Ag particles on MgO substrate
Applications of BEKP to Materials
Two Areas of application for BEKP:
Orientation Analysis
crystallographic orientation of small areas
use patterns to calculate the relationship between the crystallographic axes some external reference frame ( i.e. rolling direction)
use patterns to determine the crystallographic relationship between two adjacent areas of the sample (i.e. grain boundary misorientation, precipitate/matrix orientation)
Micro-texture imaging of polycrystalline samples with automated pattern indexing (called orientation imaging microscopy OIM)
Phase Identification
Identify unknown phases from their crystallography
Bulk samples (metallographically polished surfaces)
Particulate on substrate (no preparation needed)
Fracture surfaces (identify phases directly on fracture surfaces)
EBSD Analysis of 11 µm Diameter W wire
Microscope conditions:
20 kV, 6 nA beam current
Note the very strong <110> fiber texture developed in the wire. This is a tpical texture for BCC wire drawing.
Conventional Approaches to Phase Identification
How do we answer the “What is it?” question?
Chemistry only through EDS – may not give correct answer when phases of similar composition are present.
Combination of XRD and EDS – may work if we have sufficient amounts of material.
EBSD and EDS in the SEM – Powerful combinations of techniques with high resolution and the ability to separate compounds of similarcomposition.
How? –
Observe symmetry elements in pattern
Use angles between planes and composition to search data base
Use composition and unit cell volume to search a database
Symmetry elements in EBSD pattern
m m
m
m
4mm
m
m
m3m
m
m
2mm
Point group <111> <100> <110> <uu0> <uuw> <uvw>
m3m 3m 4mm 2mm m m 1
43m 3m 2mm m 1 m 1
432 3 4 2 1 1 1
The EBSD pattern show symmetry consistent with m3m (cubic).
Symmetry elements in EBSD pattern
Can identify 27 of the 32 point groups using this approach
Distortions of gnomonic projection can add difficulty
Generally difficult and time intensive process
Becomes more complex for lower symmetry crystal structures
Once we have the point group we can begin to infer space group
For more information see paper :
Baba-Kishi and Dingley, Backscatter Kikuchi diffraction in the SEM for identification of crystallographic point groups, Scanning, vol.11, 1989, p. 305
Phase identification using databases
We really need the qualitative chemistry to achieve phase identification
Use of EBSD pattern and chemistry:
Determine chemistry and acquire EBSD pattern
Search database based only on composition
Phases that match chemistry are then compared to EBSD patterns by simply calculating angles between planes and indexing
Use of EBSD pattern, chemistry and unit cell volume:
Determine chemistry and acquire EBSD pattern
Analyze EBSD pattern and calculate unit cell volume
Search database based on unit cell volume and chemistry
Candidate matches are used to index EBSD pattern
Phase Identification - How it is done
Obtain pattern from area of interest. Use EDS to determine chemistry.
In this case the EDS showed the presence of Fe, As and S.
Phase Identification - How it is done
Search ICDD PDF for matches based on chemistry and unit cell volume
In this example Arsenopyrite was the only match out of 100,000 compounds in the database
Phase Identification - How it is done
Index pattern based on database information.
Index pattern using the angles between the planes selected by the Hough Transform.
Phase Identification - How it is done
Simulate pattern and compare. A good match at this step is an excellent indication that the phase has been identified.
Identification of Phases on Fracture Surfaces with EBSD
.
30 µm
Austenite
Ferrite
The ocurence of hot cracks in stainless steel welds is related to the phases present in the weld. EBSD is the only technique that can directly provide this information.
Ni-Gd Phase Diagram
600
800
1000
1200
1400
1600
0 20 40 60 80 100
Wt% Nickel
1455°C
1275°C1285°C1270°C
1200°C
1110°C
1010°C
880°C
1280°C
14635°C
735°C
1313°C
1235°C
33
87
1487°C
Gd
3Ni 2
Gd
3Ni
GdN
i
GdN
i 2
GdN
i 3
Gd
2Ni 7
GdN
i 4
GdN
i 5
Gd
2Ni 17
(Ni)(αGd)
(βGd)
L
Gd Ni
Phase Identification in an Ni-Cr-Mo-W-Gd alloy
10 µm
EDS shows presence of Ni and Gd
EDS shows presence of Ni and Cr
Sample was prepared by standard metallographic techniques. No etch has been applied.
Phase Identification in an Ni-Cr-Mo-W-Gd alloy
Gd-Ni phase is identified as GdNi5(hexagonal)
Identification of Phases in a Gd modified 304 Stainless Steel Alloy
5.84 wt% Gd
Microstructures are complexThe secondary rim and the associated
gadolinides could not be identified
Sigma (FeCr)
Ferrite
Austenite
(Ni,Fe)3Gd
Ferrite
Fracture of Ta Alloy Weld
Welded component fractures along grain boundaries in the weld region.
Higher magnification shows a wide grain boundary region with precipitates. EDS shows that the grain boundary region contains only Ta and the precipitates contain Hf.
Fracture of Ta Alloy Weld (cont.)
6 possible matchesHf3N2 Trigonal (2.14% error) HfO2 Orthorhombic (6.47% error)HfO2 Monoclinic (2.25% error) HfO2 Monoclinic (1.95% error)HfO2 Tetragonal (1.64% error) HfO2 Monoclinic (1.95% error)
Fracture of Ta Alloy Weld (cont.)
Precipitate identified as monoclinic HfO2
Particle Identification - Non-volatile Memory Fabrication
Particles formed on Ru barrier layer
Particle Identification - Non-volatile Memory Fabrication
EBSP collected from facet on particle .
EDS showed that the particle contained Ru and O.
Particle Identification - Non-volatile Memory Fabrication
Pattern analyzed and crystallographic database searched. The only match that was identified was RuO2 a tetragonal phase with unit cell dimensions of 0.4499 nm and 0.6906 nm. The indexed pattern is shown at right.
Unknown Mineral IdentificationUnknown Mineral Identification
SEPatterns obtained at 20 kV Patterns obtained at 20 kV from phase containing from phase containing PbPband O. Database search and O. Database search yielded 7 possible matches. yielded 7 possible matches. Automated phase Automated phase identification selected identification selected PlattneritePlattnerite (PbO(PbO22), a ), a tetragonal phase, as the tetragonal phase, as the correct identification based correct identification based on chemistry, unit cell on chemistry, unit cell volume and pattern volume and pattern simulation. simulation.
BSE
JRM 971022
HOLZ Rings in EBSD
Mo2C (Hexagonala = 0.3012 nm c = 0.4735 nm)
Fe3C (Orthorhombica = 0.5091 nm b = 0.6743 nmc = 0.4526 nm)
Higher Order Laue Zone Analysis of BEKPs
1/λ
HG
θ
H = spacing of planes in reciprocal space
1/H = spacing of the planes in real space
H = k(1-cos 2θ) = 2k(sin2 θ)
where k = 1/λ
Hematite (Fe2O3)
5 kV 10 kV 20 kV
30 kV
kV H-1 error %
5 0.543 0.05
10 0.544 0.19
20 0.544 0.30
25 0.541 0.38
30 0.544 0.19
Actual H-1 for [211] 0.5427 nm
25 kV
HOLZ Ring Analysis of Chromium Carbide
[ 001]HOLZ
HOLZ ring around [001] used
Spacing measured normal to [001] is 1.074 nm.
The spacing of planes normal to [001] in Cr23C6is 1.067 nm, an error of 0.7%.
HOLZ Ring Measurements for SiC Polytype Identification
H-1 = 1.858 nm
H-1calc= 1.846 nm
H-1 = 1.545 nm
H-1calc= 1.538 nm
HOLZ spacing identifies left pattern as SiC 6H and right pattern as SiC 15R.
Consistent with the symmetry observed in the patterns.
Primitive Unit Cell Calculation
Any three primary vectors in the EBSP define a arbitrary unit cell representative of the Bravais lattice
If these vectors are obtained from Kikuchi lines, the cell determined is a reciprocal cell
If these vectors are determined from HOLZ ring measurements, the cell is a direct cell
Use cell reduction algorithm and the arbitrary cell to determineunique primitive reduced unit cell representative of the metric unit cell.
References:A. Santoro et al., Acta Cryst. (1980) A36, p. 796.
B. Gruber, Acta Cryst. (1989) A45, p 123.
Y. Lepage, Microscopy Research and Tech., (1992) 21,p. 158.
HOLZ Rings for Primitive Cell Calculation
Hmeas= 0.576 nmHcalc = 0.57704nm
Hmeas= 0.822 nmHcalc = 0.8089 nm
Hmeas= 0.837 nmHcalc = 0.8581 nm
Hmeas= 0.645 nmHcalc = 0.6443 nm
1
3
2
4
Measured angles between directions:
1-> 2 = 56.1° 1->3 = 44.5° 2->3 = 66.8°
HOLZ Rings for Primitive Cell Calculation
Unit cell from HOLZ Rings:
a = 0.82 nm α = 56.1°
b = 0.65 nm β = 44.5°
c = 0.65 nm γ = 66.8°
Primitive unit cell from reduction :
a = 0.58 nm α = 89.4°
b = 0.58 nm β = 112.0°
c = 0.58 nm γ = 89.0°
volume = 179 Å3
Primitive Unit cell for AsFeS :
a = 0.5741 nm α = 90°
b = 0.5668 nm β = 111.93°
c = 0.57704 nm γ = 90.0°
volume = 174.1 Å3
Primitive Unit Cell Determination of Mo 2C (hexagonal)
Zone 1H-1 = 0.48 nm
Zone 2H-1 = 0.56 nm
Zone 3H-1 = 0.71 nm
Angles between zones:
1 -> 2 = 32.0°
2 -> 3 = 24.3°
3 ->1 = 47.3°
Arbitrary Unit Cell:
a = 0.71 nm α = 32.0°
b = 0.56 nm β = 47.3°
c = 0.48 nm γ = 24.3°
Primitive Unit Cell Determination of Mo 2C (hexagonal)
Primitive Unit Cell:
a = 0.30 nm α = 88.2°
b = 0.30 nm β = 89.1°
c = 0.48 nm γ = 118.7°
Cell volume = 37.0 Å3
Published Unit Cell:
a = 0.3012 nm α = 90°
b = 0.3012 nm β = 90°
c = 0.4735 nm γ = 120°
Cell volume = 37.1 Å3
Primitive Unit Cell Determination - HOLZ Rings are not required
Primitive Unit Cell from HOLZ Rings
a = 0.27 nm α = 106°
b = 0.28 nm β = 108°
c = 0.29 nm γ = 111°
volume = 15.9 Å3
Primitive Unit Cell from Kikuchi Lines
a = 0.26 nm α = 107°
b = 0.27 nm β =107°
c = 0.27 nm γ = 111°
volume = 15.3 Å3
Primitive Unit Cell for Mo
a = 0.272 nm α = 109.47°
b = 0.272 nm β = 109.47°
c = 0.272 nm γ = 109.47°
volume = 15.6 Å3
Primitive Unit Cell Determination From Kikuchi Line Pairs
Primitive Unit Cell from Kikuchi Lines
a = 0.24 nm α = 61°
b = 0.24 nm β = 61°
c = 0.24 nm γ = 61°
volume = 10.1 Å3
Primitive Unit Cell for Ni
a = 0.248 nm α = 60°
b = 0.248 nm β = 60°
c = 0.248 nm γ = 60°
volume = 10.9 Å3
SummaryThe SEM is now a more complete tool due to the addition of EBSD to the imaging and microanalysis techniques previously available.
EBSD is a robust and reliable technique for the identification of crystalline compounds:
Use of diffraction databases allows phase identification
HOLZ ring analysis ( with proper attention to details) can be used to determine reciprocal layer spacing
Reduced cell algorithm produces recognizable primitive cells using HOLZ rings (better) or Kikuchi lines ( not as good)
These cells may be used to search existing structural databases
Application of electron backscatter diffractionto the study of phase transformations
A. F. Gourgues-Lorenzon*
The application of the electron backscatter diffraction technique to the investigation of phase
transformations is reviewed. The wide variety of results obtained using this technique is illustrated
and discussed, focusing on thermodynamics and kinetics of phase transformations, solidification,
solid state phase transformations, environmentally assisted reactions and thin film deposition.
Emphasis is also placed on two rapidly growing developments: coupling electron backscatter
diffraction with advanced experimental techniques and with more and more complex modelling of
phase transformations and of resulting material properties.
Keywords: Electron backscatter diffraction, Scanning electron microscopy, Phase transformations
List of abbreviations and symbolsACOM automated crystal orientation measurements
bcc body centred cubic (crystal structure)BKDP backscatter Kikuchi diffraction pattern
CET columnar to equiaxed transition (solidification)EBSD electron backscatter diffractionEDX energy dispersive X-ray spectrometry
fcc face centred cubic (crystal structure)FEG field emission gunFIB focused ion beamGB grain boundary
GBE grain boundary engineeringHAB high angle boundaryHAZ heat affected zone (welding)
hcp hexagonal close packed (crystal structure)HOLZ higher order Laue zone
KS Kurdjumov–Sachs (see ‘Appendix’)LAB low angle boundaryMR misorientation relationship (between product
phases)NW Nishiyama–Wassermann (see ‘Appendix’)
ODF orientation distribution function (texture)OIM orientation imaging microscopy
OR orientation relationship (between parent andproduct phases)
PTMC phenomenological theory of martensitecrystallography
SEM scanning electron microscope (or microscopy)SMA shape memory alloyTEM transmission electron microscope (or
microscopy)TRIP transformation induced plasticityWM weld metal
XRD X-ray diffraction2D two-dimensional3D three-dimensional
IntroductionThe electron backscatter diffraction (EBSD) technique,also termed automated crystal orientation measure-ments (ACOM), orientation imaging microscopy(OIM) and the Backscatter Kikuchi diffraction pattern(BKDP) technique, is utilised to determine the localcrystal structure and orientation of materials. It makesfull use of the versatility and multiscale capability of thescanning electron microscope (SEM). The principles andapplications of the technique can be found elsewhere(e.g. Schwartz, 2000; Dingley, 2004) and will be onlyshortly recalled here. Before leaving the material,backscattered electrons diffract at crystal lattice planesin the probe area. They are then intercepted by aphosphor screen (Fig. 1). Owing to Bragg diffractionconditions and to the sample to screen distance, bandscentred on diffracting planes are observed. By measuringthe location and spatial orientation of such bands onecan either determine the crystal structure or, if thecrystal structure is known, crystal orientation. Thespatial resolution is about 10–50 nm parallel to the tiltaxis, depending on the sample and operating conditions(Humphreys, 1999 and 2004; Dingley, 2004). Theangular accuracy of measurements is y1 and y0.5ufor misorientation between neighbouring areas (Dingley,2004; Humphreys, 2004) and can be improved down toy0.1u for particular applications if special care is taken(Humphreys, 2004). The technique has been developedsince the 1970s (Venables, 1973) and automatic mappinghas been readily available for about ten years. TheEBSD is used to investigate, e.g. texture and recrystalli-sation (Jazaeri et al., 2004), deformation mechanismsand local strains (Lehockey et al., 2000; Wilkinson,2000), interface characterisation (Randle, 2004) andcracking (Gourgues, 2002). Through cooperationbetween EBSD system providers and users, the techni-que has strongly developed in the past few years,including a dramatic increase in data acquisition speedthanks to digital cameras, increasing complexity of dataprocessing to extract more information from EBSD
Ecole des Mines de Paris, Centre des Materiaux PM Fourt, UMR CNRS7633, BP 87, 91003 Evry cedex, France
*Corresponding author, email [email protected]
� 2007 Institute of Materials, Minerals and Mining and ASM InternationalPublished by Maney for the Institute and ASM InternationalDOI 10.1179/174328007X160254 International Materials Reviews 2007 VOL 52 NO 2 65
data, including new pattern and map processing algo-rithms (Wright and Nowell, 2006; Brewer et al., 2006),coupling with chemical information or even reconstruct-ing parent grains as will be discussed below.
First reviews involving EBSD addressed a variety oftopics, often illustrated with a few examples (Dingleyand Randle, 1992; Mason and Adams, 1994; Randle,1994). Up to now, the number of papers mentioningEBSD is growing very rapidly (typically multiplied bytwo every year), showing that EBSD is now a routinelyused technique available in many laboratories. Thus,many reviews focused on one specific topic, such asapplicability of EBSD to ceramic materials (Farrer et al.,2000), data collection and processing (e.g. Wright et al.,2000), recent developments of the technique (Dingley,2004) and studies involving high spatial resolution(Humphreys, 2004). Only few review papers addressedthe application of EBSD to phase transformations. Thistopic is not new but it has developed only gradually. Inthe first review by Dingley and Randle (1992) only threeexamples of studies on multiphase materials were given.Applications of EBSD to solidification (Mason andAdams, 1994) and to phase identification (Randle, 1994)then began to develop. Most detailed reviews on topicsclose to phase transformations addressed phase identi-fication (Schwartz et al., 2000; Baba-Kishi, 2002). Thepresent author did not find any review paper addressingapplications of EBSD to solidification and more
generally to phase transformations in open literature,while the number of published studies is significant (togive a figure, more than 700 papers were read to writethe present review).
The present review aims to survey how EBSD is usedto address phase transformations. In this field, EBSDhas mainly been used so far to ‘revisit’ a number offeatures concerning phase transformations, the varietyof which is illustrated below, and for which crystal-lographic data were already available. This bothvalidates the EBSD technique and provides a significantnumber of data with rather simple sample preparation.The EBSD can now be used as a reference technique, yetwith limits that are discussed below, to investigate phasetransformations in the future.
Original results obtained with EBSD that were notaccessible using other techniques are also reviewed,highlighting specific advantages of EBSD in a number ofcases.
The review is organised according to transformationmechanisms, using tabular form in many cases forclarity. First, information is given on basic phenomenarelated to phase transformations (phase identification,thermodynamics and kinetics). The next two sections aredevoted to solidification and solid state phase transfor-mations respectively. Then, surface phenomena such asenvironmentally assisted reactions and thin film deposi-tion, which are of increasing practical interest, areaddressed. In the last section, coupling of EBSD withother experimental techniques and with numericalmodelling is discussed, illustrating promising ways toobtain better insight into phase transformationphenomena.
Phase identification and basicphenomena related to phasetransformations
Phase identificationElectron backscatter diffraction can be a powerful toolfor phase identification, in particular for new or complexcrystal structures or for mineral materials. The EBSD iscomplementary to longer established techniques basedon transmission electron microscopy (TEM), X-raydiffraction (XRD) and even light optical microscopy inthe case of minerals.
Crystal structure identification
This topic has been thoroughly addressed in the reviewby Baba-Kishi (2002). Here, more recent results andresults using other techniques together with EBSD arediscussed. The main advantage of EBSD versus localXRD (such as Kossel or pseudo-Kossel analysis) is itsability to perform rapid analysis, with high spatialresolution, at the expense of accuracy in crystalorientation and lattice parameter determination(Goehner and Michael, 1995; Dabritz et al., 2001).Lattice parameters can be evaluated either from theband width (which is proportional to the sine of thecorresponding Bragg angle) or from the size of higherorder Laue zone (HOLZ) rings in EBSD patterns (Baba-Kishi, 2002). The accuracy of measurements depends ondata acquisition conditions and on the chemicalcomposition of the analysed area. By using HOLZ ringsit can be better than 1%, if care is taken to correct for
(a)
(b)
a Electron diffraction according to Bragg’s law(2dhklsinh5nl); b EBSD pattern from a bainitic steel (bcccrystal structure)
1 Principle of EBSD analysis
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
66 International Materials Reviews 2007 VOL 52 NO 2
distortions due to the gnomonic projection or to atomicnumber (Baba-Kishi, 1998; Michael, 1999 and 2000;Michael and Eades, 2000). However, for elements suchas Mo or W, the error is still about a few per cents(Michael and Eades, 2000). The location of HOLZ ringsin EBSD patterns can also be used to distinguishbetween polytypes, in particular in clay minerals(Kogure et al., 2005). Using the band width gives loweraccuracy [typically a few per cents (Michael andGoehner, 1993)] even with increasing the specimen toscreen distance or optimising the indexing algorithm(Vatne et al., 1998; Dingley, 2004). Extinction effects(Michael and Eades, 2000) and cell multiplicity (Vaillantet al., 2003) must be taken into account. The feasibilityof EBSD assisted phase identification (e.g. Michael andGoehner, 1993) essentially depends on the ‘signal tobackground noise’ ratio (Small et al., 2002), on back-ground acquisition (Michael, 2000; Small and Michael,2001) and on relief effects (Michael, 2000), not much onthe yield of backscattered electrons (Small and Michael,2002). Therefore, good results can be obtained withextraction replicas [e.g. for PbO2 (Michael, 2000) andU3O8 (Small and Michael, 2001)]. Particles over about0.3–1 mm in size can be readily analysed with EBSD(Straumal et al., 1999; Small and Michael, 2001).
Crystal symmetry has been addressed with EBSD fora long time (Baba-Kishi, 2002). 27 out of the 32 pointgroups can be identified using EBSD patterns.Identification methods of point and space groups aredetailed elsewhere (Baba-Kishi and Dingley, 1989;Michael 2000; Baba-Kishi, 2002). This field has stillbeen developed since these review papers [(see e.g. Kralet al., 2004) for iron rich phases in an Al–Si alloy], inparticular for minerals and quasicrystals. Super-structures of ordered intermetallic phases such as theL10 c phase in TiAl alloys are difficult to investigate byrapid, automated analysis, so that they are often ignoredif the distinction between them is not of utmost impor-tance for the considered application (Pouchou et al.,2004b). Special attention has been paid to minerals formany years (e.g. Dingley, 1984; Leinum et al., 2004) andin particular to polytype identification. Kogure (2003)developed a method based on band location andintensity, still using kinematic theory of electron–matterinteractions, which shows recognisable patterns char-acteristic of each polytype. This method gives goodresults for phyllosilicate minerals (Kogure, 2002 and2003; Kogure and Bunno, 2004; Kogure et al., 2005).Others established robust reflection tables (i.e. extinctionconditions) as e.g. for omphacite (Mauler et al., 1998).Identification of quasicrystal structures still generallyrequires manual indexation of patterns [e.g. forAl60Cu26Fe14 (Cheung et al., 2001)]. Special care mustbe taken to check the symmetry (including from insidethe bands) around zone axes, to distinguish ‘true’quasicrystals from their crystalline approximants(Ruhnow et al., 2002).
Coupling EBSD with chemical analysis
Energy dispersive X-ray spectrometry (EDX) microana-lysis is frequently available in the SEM together with theEBSD facility, allowing combination of chemical andcrystallographic information (although optimised oper-ating conditions are not the same). Starting fromsequential analysis of the same area, EBSD can be usedto facilitate EDX-based phase identification, especially
for off stoichiometry phases and for phases havingidentical or similar chemical compositions but differentcrystal structures (Camus, 2000; Kogure, 2002;Boettinger et al., 2003; Kogure and Bunno, 2004; Kralet al., 2004; Leinum et al., 2004). Up to date availablesystems now combine both crystallographic and chemi-cal analysis for automated phase identification (Michael,2000; Dingley, 2004; Nowell and Wright, 2004; Zhonget al., 2006). For rapid data acquisition and processing,the location of bands in the EBSD pattern is stored foreach data point, together with chemical analysis. Thechemical composition is then used to index every patternin an offline manner. This is especially useful when theidentity of phases is not known in advance, or if datacollection must be rapid (e.g. for in situ investigations).The spatial resolution is limited, however, to that ofEDX analysis (y1 mm), both for in plane dimensionsand along the sample normal.
Distinguishing between known phases
Minerals have been studied with EBSD due to theirfrequently low crystal symmetry and to the widecollection angle (y90u) provided by this technique. Insome instances, a list of reflections needs to be firstcalculated or corrected and then given to the EBSDsystem [e.g. omphacite in eclogite minerals (Mauler et al.,1998) and Al–Si–O triclinic phase in kyanite (Dingleyet al., 2004)]. In some cases manual adjustment is stillnecessary to insure right indexation of the EBSDpatterns [e.g. garnets and sulphides (Prior et al.,1999)]. While researchers in the past had to work withfilms to improve sensitivity and image contrast (Dingley,1984), new digital cameras now allow online working inthe SEM. The case of polytype minerals involves bothnumerical calculation of the patterns and visual identi-fication of individual patterns, provided that crystals aresuitably oriented (Kogure, 2002 and 2003; Kogure et al.,2002 and 2005; Kogure and Bunno, 2004; Kameda et al.,2005). In many cases, technical improvements of EBSDcameras allow switching from manual to automatedEBSD [e.g. for iron ore and sinter (Magalhaes, 2002;Sasaki et al., 2005a)]. Advantages and challenges ofusing EBSD to investigate ceramic materials havealready been reviewed (Farrer et al., 2000) and will notbe recalled here. Recent studies addressed an increasingvariety of materials such as oxides (Michael andGoehner, 1993; Rogers et al., 1994; Randle, 1994;Korte et al., 2000; Cha et al., 2006), functional materials[YBCO (Koblischka et al., 2003; Koblischka-Venevaet al., 2003)], zeolites (Pennock et al., 2001), PbTiO3
domains (Yang et al., 1994), silicon (Her et al., 2000),eutectic fibres (Nakai et al., 2005; Lee et al., 2005),materials for nuclear applications (Medevielle et al.,1999) and applications are still rapidly growing. As anillustration, example studies of quasicrystals and metal-lic glasses are reported in Table 1.
Phase identification is also useful to understand in situsynthesis of composites through interfacial reactions.Good examples are given by AlN/Ti composites leadingto several (Ti,Al)Ny nitrides (Paransky et al., 1999,2000a and 2000b) and by sintered WC/VC/TiC/Cocermets (Arenas et al., 2005). Alignment of strengthen-ing particles during composite processing (Schuh andDunand, 2001), formation of intermediate phases duringtransient liquid phase bonding of a Ni base superalloy(Jalilian et al., 2006) and carbide formation in SiC/Ti
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 67
alloy filamentary composites (Duda and Gourgues,2006) were also investigated with EBSD.
As SEM is much easier on electron conductingsamples, most EBSD studies have yet been carried outon metals and metal alloys. Besides many others, somespecifically address phase identification, as for exampleidentification of small precipitates [e.g. NiSn4
(Boettinger et al., 2003), MX particles in Ni base alloys(Ramirez and Lippold, 2004), discontinuous precipitatesin ancient components (Wanhill, 2005) and Laves phasesin austenitic alloys (Michael and Goehner, 1994; Robinoet al., 1997)], of topologically close packed precipitatesin tempered weldments (El-Dasher and Torres, 2004), ofeutectic aggregates in a Ni base alloy (DuPont et al.,1999) and of complex/multilayered aggregates of phases(Levin et al., 2000; Gomez and Echeberria, 2003; Nolzeet al., 2005). First principles calculations have been usedto find crystal structures and lattice parameters forfurther automatic indexing of EBSD patterns [e.g. innewly discovered Al2(Mg,Ca) ternary Laves phases(Zhong et al., 2006)]. The case of alloys particularlydifficult to investigate such as Pu–Ga alloys (Boehlertet al., 2001) and cerium (Boehlert et al., 2003a) should bementioned. Analysis of EBSD also allows convenientphase separation on fracture surfaces (not readilyavailable with TEM) (Foct and Akdut, 1993; Bacheet al., 1997a and 1997b) and for texture determination(not always easy by XRD) (Xu and Russell, 2004; Dunstand Mecking, 1996). The EBSD can also facilitateevidence of phase transformation (Sano et al., 2003b;Neishi et al., 2004; Zhu, 2004).
Discrimination by EBSD between phases having closeor even identical crystal structures is difficult. In somecases, the lattice parameter is used [e.g. to separate g9T
from e in Zn–Al–Cu alloys (Zhu et al., 2001, 2003a and2003b)]. In other cases, the quality of the patterns ismuch better in one phase than in the other [e.g. ferrite v.martensite (Wilson et al., 2001; Wilson and Spanos,2001; Jeong et al., 2002a and 2002c; Petrov et al., 2004;Cabus et al., 2004b; Wu et al., 2005), and ferrite v.bainite (Regle et al., 2004) in steels]. Primary andsecondary a phases in a titanium alloy (Germain et al.,2005a) were distinguished using backscattered electron
imaging. An alternative method using light opticalmicroscopy was also developed (Thomas et al., 2005).Discrimination between such phases also allows map-ping spatial and size distribution of them (see e.g.Schwarzer et al., 2000).
Equilibrium phase diagram assessmentRapid phase identification by EBSD, together withassessment of properties of individual constituents bylocal measurements has recently been reviewed in theframework of diffusion multiples (Zhao, 2006). Withthis technique, a whole isothermal section of one (orseveral) ternary diagram can be assessed after only onelong term heat treatment. This has been applied tovarious ternary alloys, assuming local equilibrium atinterfaces. However, the effect of local crystallographyon diffusion and on phase nucleation was not discussed(Zhao et al., 2001, 2002, 2003, 2004a, 2004b; Zhao, 2004;Zhong et al., 2006). Modification of existing phasediagrams can be suggested by EBSD results [e.g. for1200 and 1600uC isothermal sections of the quaternaryMo–Ti–Si–B diagram (Yang et al., 2005)]. Applicationof EBSD to local adjustment of the liquidus surface ofthe Ti–Al–(Nb) system is illustrated by identification ofthe solidification sequence (Takeyama et al., 2004).Except for this latter case, EBSD was mostly used as acomplementary technique to discriminate betweenphases having complex and close chemical compositionsbut different crystal structures. Rapid mapping facilitiesthat are now available should put EDX plus EBSDas leading experimental techniques for phase diagraminvestigations in the next future.
Electron backscatter diffraction and kineticissues of phase transformationsAnisotropy of ‘free’ surface and interfacial energy
‘Free’ surface energy
The equilibrium shape of single crystals is a goodtranslation of free surface energy into a measurableproperty. Provided that crystals are suitably orientedwith respect to the SEM, surface planes can be indexedwith EBSD, as e.g. ZnO well faceted fibres (Huang et al.,2004), clay minerals (Kogure et al., 2005; Kameda et al.,
Table 1 Phase identification assisted by EBSD in quasicrystals and metallic glasses
Material Phases of interestPhase identificationmethod Results Ref.
Cast Al–Cu–Fe alloys b-Al(Cu,Fe) (CsCl) Cu rich phases,e phases, icosahedral quasicrystallineand R phases, l-Al13Fe4 dendrites
EBSD z EDX Crystal orientations,orientation relationshipsbetween phases
Gui et al.,2001
Cast Al–20Cu–15Fe(at.-%)
Al60Cu26Fe14 (icosahedral), b-Al(Cu,Fe)(CsCl), near-l Al44Cu54Fe2
Manual indexing anduse of ‘Kikuchi’ atlas
Structural model of theicosahedral phaseconsistent with TEMresults
Cheung et al.,2001
AA6013 Al alloy Fe rich phase Symmetry in EBSDpatterns, comparisonwith XRD results
Not a quasicrystallinephase but crystallineapproximant
Ruhnow et al.,2002
Al–Cu–Fe–B b-Al(Cu,Fe) (CsCl), Fe2AlB2 (bodycentred orthorhombic), quasicrystallinematrix
Point analysis Identification of phases,orientation relationshipsbetween phases
Brien et al.,2004
Fe–Co–Zr–Mo–W–Bmetallic glass
ZrB2 (hexagonal) Automatic indexing withvarious crystal structuresand lattice parameters
Identification of thisresidual crystallinephase
Castellero et al.,2004
Zr–Cu–Al–Ni metallicglass
Zr–Cu–Al–Ni–O [face centred cubic(fcc)]
Use of band width Identification of phasestemming from oxygencontamination
Vaillant et al.,2003
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68 International Materials Reviews 2007 VOL 52 NO 2
2005) and tape cast ceramics (Markondeya Raj et al.,1999). Surface energy of SrTiO3 perovskite has beeninvestigated by trace or 3D plane analysis after boththermal grooving (Sano et al., 2003a) and coarseningwith TiO2 rich liquid (Sano et al., 2005). Interfacemigration assisted by an amorphous phase has beenstudied in sapphire single crystals bonded to sapphirepolycrystals as a function of the local orientation of‘free’ surface planes (Farrer et al., 2006). Relativeenergies of low index planes can be rather easilycompared by this way with minimum sample prepara-tion and on a larger data basis compared with TEMtechniques. On the other hand, assumptions have to bemade to accurately determine the spatial orientation ofthe free surface planes (see ‘Discussion’).
Internal interfaces
Interfaces are good candidates for heterogeneous phasenucleation, depending on their free energy. Electronbackscatter diffraction has been used in many ways inthis field. Interface wetting [e.g. Zn by Ga (Traskineet al., 2005), partially molten olivine aggregates (Fauland Fitz Gerald, 1999)] involve classification of ‘dry’and ‘wet’ interfaces according to e.g. the misorientationangle between neighbouring crystals (this being readilyaccessible with EBSD). The effect of lattice coincidencein these two cases is not obvious. The same result wasobtained with solid state spreading of copper on coppersubstrate (Missiaen et al., 2005). Thermal groovingcoupled with EBSD and near field microscopy givesaccess to interface energy anisotropy. Here again,coincidence site lattice criteria fail to qualitativelycompare values of various grain boundary energies[e.g. in MgO (Saylor and Rohrer, 1999) and NiAl(Amouyal et al., 2005)]. In the latter case, the boundaryenergy correlates with none of usual criteria such ascoincidence site lattice, misorientation angle, tilt v. twistcharacter or plane boundary matching.
Segregation of solute atoms to boundaries has alsobeen investigated with EBSD in hydrogen containingAl–5Mg alloy (Horikawa and Yoshida, 2004) and Nbdoped rutile (TiO2) (Pang and Wynblatt, 2006).Coincidence site lattice boundaries are often resistantto solute atom segregation and thus to localisedcorrosion and fracture [e.g. in Ni (Cornen and LeGall, 2004), in boron containing AISI 304 stainless steel(Kurban et al., 2006) and for intergranular fracture ofFe–0.002C–0.06P (Williams et al., 2004)]. By readilyacquiring a high number of data, EBSD allows study ofthe ‘grain boundary character distribution’, which is inturn influenced by the interface energy distribution, as afunction of parameters such as chemical composition. InSn–Ag–Cu lead free solder alloys, Telang et al. (2004)showed that Ag (respectively Cu) tends to reduce(respectively increase) the difference in free energybetween various types of boundaries in the material.
Diffusion along interfaces
Interface diffusion may be characterised with the help ofEBSD as a function of the local interface structure. Themain difficulty is accurate determination of the localcrystallographic orientation of the interface plane(Randle, 2004). Discontinuous ordering of Fe–50 at.-%Co highly depends on the grain boundary(GB) misorientation. ‘Special’ GBs such as low angleboundaries (LABs) and S3 twin boundaries, allowing
only slow diffusion, are not sensitive to discontinuousordering. Vacancy migration is very difficult if no,111. direction is close to the GB plane, so that thesensitivity of high angle boundaries (HABs) to discon-tinuous ordering depends on both GB misorientationand local orientation of the GB plane (Bischoff et al.,1998; Semenov et al., 1998). Surface rearrangement inFeO is also strongly affected by neighbouring GBs,whatever the local free surface orientation (determinedby EBSD), due to preferential GB diffusion (Bahgatet al., 2005). Depending on the local crystallographicorientation of free surfaces, a Kirkendall effect due todifferences in solute atom diffusion can be observedafter phase trans formation [e.g. in Ti–9 wt-%Mo (Guoand Enomoto, 2006)].
Anisotropy of bulk diffusion
Anisotropy of bulk diffusion strongly depends on thecrystal structure. The EBSD was used to investigatediffusion of Cr in tetragonal Mo5Si3, for which thediffusivity along [100] and [010] directions is higher thanalong the [001] direction (Strom and Zhang, 2005).When single crystals are not readily available, diffusionin coarse grained materials after EBSD identification ofeach grain orientation is a convenient tool in this field[see e.g. oxygen tracer diffusion in (La2–xSrx)CuO4
(Claus et al., 1994)].
Solidification and semisolid state
Nucleation of solid phaseNucleation occurs at a very local scale, which cannot beaccessed with the SEM. However, one of its conse-quences (i.e. crystal orientation) can be efficientlystudied with EBSD, in particular when sample prepara-tion is tedious (e.g. for composite or heterogeneousmaterials such as welded joints), or if the microstructureis too coarse to be characterised with TEM or XRD. Byallowing access to local crystallography over large areas,the EBSD technique facilitates investigation of phenom-ena that are difficult to access by TEM, such asheterogeneous nucleation at a substrate surface or atscarcely distributed particles. Examples are numerousand will still grow rapidly in number in the next future.
Effect of nucleating agents
A few EBSD studies address nucleation enhancement(or inhibition) by particles or solute elements. Specificorientation relationships (ORs) between nucleatingparticles and solidifying material have been observedin, e.g. aluminium inoculated with a TiC forming alloy[cube–cube OR (Tronche and Greer, 2001)] and in AISI409 ferritic stainless steel at TiN particles ({001}TiN//{001}ferrite and ,110.TiN//,100.ferrite) (Hunter andFerry, 2002a) leading to highly textured material.Equiaxed, randomly textured grains nucleate at segre-gated areas of the heat affected zone (HAZ) of Sc or Zrcontaining aluminium alloys (Kostrivas and Lippold,2004). Sb and Pb were shown to strongly inhibitnucleation of the solid phase in zinc based hot dipcoatings (Quiroga et al., 2004; Rappaz et al., 2004).
Solidification path
The EBSD has been intensively used for stainless steels,whose solidification path strongly depends on bothchemical composition and solidification rate. Theaustenite (fcc c) phase generally forms first in rapid
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International Materials Reviews 2007 VOL 52 NO 2 69
solidification, whereas ferrite (bcc d) forms first at lowersolidification rates. This has been illustrated in laser(Robino et al., 1998; Katayama et al., 1999; Brookset al., 2003; Iamboliev et al., 2003) and gas tungsten arc(Robino et al., 1998; Iamboliev et al., 2003) weldedalloys. As d and c phases exhibit particular ORs whenone forms by solid state transformation from the other,it is rather easy to determine by EBSD, with informationon both phase morphology and crystallography, whichphase was the first to form, if any. When a strong textureis observed, but not these particular ORs, one cansuppose that both d and c formed simultaneously duringsolidification. This could be the case in weld metal (WM)deposits (Bouche et al., 2000) (Fig. 2). The EBSDprovides information, which in these coarse grainedmaterials cannot by obtained with XRD, and on a muchmore reliable statistical basis than by TEM.
Information about abrupt changes in solidificationmode can be given by EBSD. By continuously varyingCO2 laser welding speed or base metal composition ofFe–18Cr–(10–14)Ni alloys, Fukumoto and Kurz (1997)showed that the change in solidification mode occurredwith no change in crystal orientation and no nucleationof new misoriented crystal, but after some undercooling.Owing to the higher distance between d and the cdendrite tips than between c and the d dendrite tips, theFA (d then c) to AF (c then d) transition occurred morereadily than the AF to FA transition.
The EBSD proved useful to investigate the solidifica-tion path of MCrAlY alloys deposited on a Ni basesuperalloy by continuous CO2 laser cladding (Bezenconet al., 2003), as soon as the lattice misfit could beaccommodated by LABs, the fcc c phase first formed byepitaxial growth. As soon as another primary phaseformed (e.g. b-NiAl), nucleation occurred and epitaxywas no longer possible. The solidification structure ofsmall areas in heterogeneous materials was also inves-tigated by EBSD at the right scale in laser deposited Ti–10 at.-%Cr alloys (Banerjee et al., 2002), in eutecticareas of Gd containing Ni–Cr–Mo alloy ingot (Robinoet al., 2003) and in Mg–Li–Ca cast alloys (Song andKral, 2005).
A recent example of application of EBSD to asolidification problem of great practical significancewas given by Sengupta et al. (2006), who studiedoscillation marks and hooks in continuously cast lowcarbon steel slabs. As the meniscus line still remained aHAB, it was clearly visible even after several phasetransformations occurred during cooling. It wasdeduced that solid phases did not nucleate at the sametime on both sides of the frozen meniscus. Owing to thelarge grain size and sample geometry, such informationcould only be accessed using EBSD.
Eutectic solidification
Nucleation of coarse eutectic grains was first studiedwith EBSD in white cast irons (Randle and Laird II,1993). A colony of hexagonal M7C3 carbides wassuggested to originate from a single nucleus. Most otherstudies addressed alloys undergoing dendritic, theneutectic solidification, to check whether the crystalorientation of the eutectic phase was the same as thatof its dendritic counterpart, i.e. if there was or notnucleation of that phase from the undercooled liquid. Inan arc melted hypereutectic Nb–Si alloy, Drawin et al.(2005) showed that Nb3Si shared a common crystal
orientation in eutectic colonies and in some dendriteneighbours. The vast majority of published workaddressed the solidification of Al–Si alloys modifiedwith Sr, Sb, Cr, Na and P and cast as small ingots. Thechange in morphology of eutectic silicon from lamellarto fibrous could not be unambiguously related to achange in OR between dendritic and eutectic (Al) phases(Dahle et al., 2001). While in unmodified alloys eutectic(Al) had most often the same crystal orientation as aneighbouring dendrite, this was not the case for
(a)
(b)
2 a light optical micrograph of type 316L stainless steel
WM and b {110} pole figure, in c austenite frame, of d
ferrite (individual dots) located at c LABs, showing
either KS ORs (surrounded by circles and triangles) or
near cube–cube ORs (surrounded by rectangles)
between d and c. After Bouche et al. (2000)
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70 International Materials Reviews 2007 VOL 52 NO 2
modified alloys (Wang et al., 1999; Nogita and Dahle,2001a, 2001b and 2001c; Nogita et al., 2001; Dahle et al.,2005), except for high amounts of added Sr (Dahle et al.,2001; Nogita and Dahle, 2001a). Strontium was thoughtto increase the efficiency of nucleating agents, leading toindependent nucleation of eutectic (Al) from the liquid,just as for a columnar to equiaxed transition (CET)(Nogita and Dahle, 2001a). In contrast with thesecommercial purity alloys, Heiberg and Arnberg (2001)and Heiberg et al. (2002) did not find any CET in highpurity Al–Si–(Sr) alloys; eutectic and dendritic (Al)phases shared the same crystal orientation and thus nobarrier to nucleation of eutectic (Al) from dendritic (Al)was found. Electron backscatter diffraction on both (Al)and (Si) phases (Heiberg and Arnberg, 2001) showed noparticular change in internal twinning of eutectic (Si)and no particular OR between both phases, althougheutectic (Si) had a ,110. fibre texture. Here again, theneed to characterise fine scale microstructures withincoarse grains obviously suggests EBSD as the mostsuitable technique to get statistically reliable andmetallurgically relevant information.
Peritectic solidification
Rather few studies have been dedicated to peritecticsolidification except for stainless steels (see section onsolidification path). In cast iron, no OR was foundbetween newly formed M3C and already existing M7C3
carbides at which they very likely nucleated (Randle andLaird II, 1993). The peritectic c2c9 solidification ofmodel and CMSX–4 single crystal nickel base super-alloys (Warnken et al., 2005) was studied at coarseshrinkage cavities; ball shaped c9 particles had a cube–cube OR with the dendritic c substrate.
Orientation relationships between seed (or substrate ormould) and solid
Chill or external zones of ingots or thin strip castingshave frequently been shown to exhibit a fine grained,randomly textured microstructure, except for some caseswhere a strong fibre texture was already observed with
EBSD (Summers et al., 2004). At a very local scale, arandom texture can be induced by nucleation at a roughsubstrate followed by locally textured growth [e.g. instainless steels (Hunter and Ferry, 2002b; Ferry andHunter, 2002)]. Numerous misoriented equiaxed grainswere also observed in molten silicon droplets solidifiedon a silicon wafer substrate (Nagashio et al., 2004)although liquid spreading did not allow columnargrowth perpendicular to the contact plane in thatparticular case.
Single crystal superalloys such as CMSX–4 andCM186LC exhibit a randomly textured chill zone.With EBSD one can check that the distribution ofmisorientation between grains follows the classical curvefor textureless cubic materials (Ardakani et al., 2000).
Epitaxy between the HAZ and the WM can be readilycharacterised with EBSD at the local (grain to grain)scale (Table 2). The EBSD was also used to evaluateORs between seed and solidified material in complexmicrostructures, as in fully lamellar c-TiAl alloys. Here,various grains appeared at the seed, sharing one ,110.
direction with it. In Ti–48Al the selected grain had thesame crystal orientation as the seed. In Ti–48Al–8Nb afew variants having one ,110] direction of lamellaeparallel to one of the seed were observed (Takeyamaet al., 2004).
Functionally graded materials may be obtained bylaser melt injection. Interfacial reactions and ORsbetween the injected ceramic powder particles and thehot metal matrix can be characterised with EBSD ata very local scale, allowing better insight into thesolidification mechanism. Ocelık et al. (2001) examinedone hundred interfaces between Al4C3 and SiC powderin the reaction layer of laser melt injected Al–SiCcomposites. In 25% of observed cases, {0001}SiC wasparallel to {0001}Al4C3
, according to the angle betweenthe {0001}Al4C3
plane and the temperature gradient. Incontrast with extrusion (where Al4C3 is not formed), noOR was observed between SiC particles and (Al) matrix.Specific ORs were observed in reaction zones of
Table 2 Orientation relationships found with EBSD in weld solidification
Base metal Weld metal OR between WM and HAZ Ref.
Be alloy Autogeneous electronbeam welding
Epitaxy at HAZ equiaxed grains Wright and Cotton,1995
Al alloys GTAW* with or withoutfiller metal
Equiaxed WM: no OR, no texture;dendritic WM: epitaxy on HAZ
Kostrivas and Lippold,2004
Gd containing Ni base alloy Autogeneous electronbeam and GTAW
Columnar WM in epitaxy with HAZ Robino et al., 2003
CMSX–4 Ni base superalloyzMCrAlY Laser cladding Cube–cube OR if primary solidificationinto c; no OR otherwise (b or b z c)
Bezencon et al., 2003
AISI 304 and 310S austenitic stainlesssteels
Pulsed YAG laser Cube–cube OR for solidification into c,no clear OR for solidification into d
Katayama et al., 1999,Iamboliev et al., 2003
AISI 304 austenitic stainless steel GTAW Kurdjumov–Sachs (KS) OR betweenferritic WM and austenitic HAZ
Iamboliev et al., 2003
Free machining stainless steel Autogeneous pulsedYAG laser
Epitaxy (no fusion boundary) if solidifiedinto c; no clear OR if solidified into d
Brooks et al., 2003
High purity Fe and Monel 70–30 filler GTAW Nishiyama–Wassermann (NW) and BainORs{
Nelson et al., 1999a
2.25Cr–1Mo steel and 625 Ni base filler GTAW Mostly Bain OR{ Nelson et al., 1999aA508 low alloy steel and AISI 309Laustenitic filler
GTAW Bain, sometimes NW OR Nelson et al., 1999a
AISI 409 ferritic stainless steel andMonel 70–30 filler
GTAW No OR (chill zone) Nelson et al., 1999b
*GTAW: gas tungsten arc welding.{In fact, cube–cube OR during solidification into austenite and then solid state transformation into ferrite in HAZ.
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functionally graded Ti–6Al–4V/SiC or WC compositesproduced by laser melt injection, depending on the facetsof powder particles (Pei et al., 2002; de Hosson andOcelık, 2003). A functionally graded aluminium alloycontaining Al3Ti particles showed no OR betweenphases in centrifugally casting, leading to a morpholo-gical texture of Al3Ti platelets with no crystal texture inthe (Al) matrix (Watanabe et al., 2002).
Growth of solid phasesGrowth direction
Studies of EBSD on growth direction with respect to thethermal gradient have concentrated on directionalsolidification, where microstructures are frequentlycoarse grained (Table 3). In eutectic and peritecticmicrostructures, morphology and crystallographic ORsare intimately related, as reported below (section onmicrotextures). As an example of dendritic solidifica-tion, Nagashio and Kuribayashi (2005) associatedEBSD with light confocal microscopy (Fig. 3) toevidence the transition in growth direction from,211. to ,110. and then to ,100. for splatquenching of silicon onto silicon wafers. The transitionwas interpreted with facet planes and liquid/solid interfacial energy thanks to careful experimentalobservations.
Grain size and morphology
The EBSD technique is a good means to evaluate the grainsize and shape on a quantitative basis, grains being definedby boundaries of misorientation angle higher than a userfixed threshold. This implies that the ‘cleaning’ procedureused for EBSD data processing should be suitably chosenand clearly reported. Some examples in as solidifiedmicrostructures are given in Table 4. The EBSD isparticularly useful when the grain size is similar to thesize of the part, i.e. for so called ‘multicrystals’. This isthe case, for example, for lead free, coarse grained singleshear lap Sn–3.5Ag solder joints (Telang and Bieler, 2002,2005a and 2005b) and for Co–Cr–Mo laser cladding onrailway wheels (Farooq et al., 2006).
The EBSD may also reveal that individual particlesare in fact single crystals [e.g. coarse Bi particles in Bi–24.8In–18.0Sn eutectics (Ruggiero and Rutter, 1995)]or polycrystals composed of several grains [e.g. M7C3
carbides in high chromium white cast iron (Powell andRandle, 1997)].
Interrupted solidification tests allow investigatingmicrostructure formation, although the quenched liquidmay be difficult to distinguish with EBSD from thegrowing solid. The EDX may be here of great help tostudy solid/solid grain boundaries already formed beforequenching (Liu et al., 2005). In other cases, EBSDpatterns from the quenched liquid cannot be indexed, sothat the shape of the solid/liquid interface is easilyrevealed with this technique [e.g. in Fe–4.3Ni (Farynaet al., 2002)].
Phase connectivity is difficult to assess from lightoptical or electron microscopy imaging only. Byidentifying areas of the material sharing a commoncrystal orientation, EBSD provides interesting informa-tion even from two-dimensional (2D) sections, i.e.individual crystals of a 2D section may in fact belongto the same three-dimensional (3D) grain. Someexamples are given in Table 5. This is particularly usefulwhen individual grains are connected in 3D, but notin 2D (e.g. in EBSD maps or thin foil TEM images).Another side of the problem is grain clustering:misorientation gradients, which may influence productproperties, are readily revealed with EBSD as long as
3 Results of EBSD on determination of local dendrite
crystallography in splat quenched silicon: after
Nagashio and Kuribayashi (2005)
Table 3 Results of EBSD on solidification growth direction
Material Solidification process Solidification mode Crystal direction of solid growth Ref.
Ti–26Al–27Nb–0.03O(at.-%)
Induction float zonemelting
Directional solidification ,100.bcc Boehlert and Bingert,2001
Pb(Mg1/3Nb2/3)O3 z
PbTiO3
High pressureBridgman
Cellular [110]rhombohedral, sometimes 12–20ufrom it (trace analysis)
Soundararajan et al.,2004
Nb–33Ti–16Si (at.-%) Czochralski Cellular–dendritic [113]Nb; [001](Nb,Ti)3Si Sutliff and Bewlay,1996
AZ91 Mg alloy Bridgman Columnar ,1120. or ,2445. according tothermal gradient and growth rate;can be primary and ternary ,1120.
and secondary (,1120. z ,2445.)
Pettersen and Ryum,1989; Pettersen et al.,1990
Be alloy Electron beam welding Columnar ,1010. favoured Wright and Cotton,1995
Fe–3Si Twin roll casting Columnar Often close to ,100.bcc Takatani et al., 2000Al–Zn–Si Hot dip galvanising Dendritic–columnar
(Al)Neither ,100.Al nor ,110.Al,close to ,320.Al (trace analysis)
Semoroz et al., 2001
Zn Hot dip galvanising Dendritic (Zn) Mainly ,1010.hcp Semoroz et al., 2002bTi–6Al–4V Vacuum arc remelting Columnar ,100.bcc (leading to ,1120.hcp)
0 and , 45u from heat flowGlavicic et al., 2003c
Ni base Alloy 22 Gas tungsten arcwelding
Dendritic–columnar ,100.fcc (followed by partialrecrystallisation)
El-Dasher et al., 2006
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72 International Materials Reviews 2007 VOL 52 NO 2
mapping is possible. Grain clustering may be related tothe solidification texture, which increases the frequencyof LABs (West and Adams, 1997). Here, grain size andgrain connectivity strongly depend on the definition ofgrains given by the EBSD user.
Competitive grain growth
During crystal growth, a strong selection may occuramong grains nucleated either in the chill zone or at free
surfaces [e.g. in laser metal forming (Gaumann et al.,1999)] or at the mould/metal interface or even from theundercooled melt. This competition is governed by thegrowth rate of individual grains, which is in turn relatedto the local crystal orientation with respect to the localheat flow.
During directional solidification, competitive graingrowth leading to a columnar zone with an increasinggrain size (e.g. Fig. 4) was clearly evidenced thanks to
Table 4 Electron backscatter diffraction characterisation of grain size and morphology of as solidified microstructures
Material Solidification process Grain size or morphology Ref.
2024 Al alloy Metal inert gas welding More than 80% of the WM 5
equiaxed grains of 50–100 mmin size
Lefebvre et al., 2005
Si Simulated melt spinning Equiaxed grains nucleate in thechill zone, followed by columnargrowth
Nagashio and Kuribayashi,2006
Si Splat cooling on Si wafer Many equiaxed grains near thecentre
Nagashio et al., 2004
AISI 316L stainless steel Autogenous CO2 laserwelding
Grain size 200 mm; cellular (size 3 mm)if the laser beam is focused, non-cellularif irradiation is uniform
Kell et al., 2005
2618 Al alloy Thixoforming Globular grains with HABs if cast atliquidus temperature
Xia and Tausig, 1998
Inconel 82 Ni base alloy Tungsten inert gas weldingwith low carbon steel
Grain size about 150–300 mm dependingon magnetic stirring
Kokawa et al., 1999
5052 and 5182 Al alloys Direct chill and twin rollcasting
Same grain size found with EBSD andlight microscopy (but not in the sameobservation plane)
Slamova et al., 2003
Al–0.15Fe–0.07Si Direct chill casting Same grain size (140 mm) found withEBSD and light microscopy
Samajdar and Doherty,1994
AZ91D and AM60BMg–Al alloys
High pressure die casting Grain size distribution before solutionannealing (not possible in light microscopy)
Bowles et al., 2004
Fe–15 at.-%Ga Roll casting Grain size 6.3–7.9 mm Saito et al., 2004Al–7Si–0.3Mg As cast with artificial
shrinkage porosityAverage grain size ,300 mm Buffiere et al., 2001
Al–7Si–0.4Mg Low temperature semiliquiddie casting
Cells clustered into single crystal coloniesdelimited by HABs
Han et al., 2001
Si–30Al Spray forming Grain size (y5 mm for Si and 250 mm for Al)in the divorced eutectic formed after slowcooling; the difference is due to nucleationconditions (Si from droplets, Al within theformed material)
Hogg et al., 2006
Table 5 Phase and grain connectivity determined with EBSD in as solidified microstructures
Material Solidification process Results obtained with EBSD Ref.
Al–1.25Mn Directional (cellular) Highly interpenetrated grains Sun and Ryum, 1992White cast irons As cast Hypereutectic alloy: M7C3 carbides
are single crystalline over large areas(shown by deep etching to beinterconnected); hypoeutectic alloy:M7C3 carbides are individual crystallites
Randle and Powell, 1993
High Cr white cast irons As cast Connectivity of carbides varies from onematerial to the other
Powell and Randle, 1997
Sn–17 wt-%Pb Quenched from semisolidstate
2D spatial connectivity is not correlatedto crystal orientation: 3D percolation issuggested; no coalescence (density ofLABs decreases with increasing time):Ostwald ripening
Wolfsdorf-Brenner et al., 1999
Sn–1.4 wt-%Cd Directionally solidifiedoscillating peritectic
a phase: multilayered grains connectedin 3D; b phase: connected even in 2D
Zeisler-Mashl et al., 1997
Ni–Cr–Co–Ti–Al–Mo–Sisuperalloy
Directionally solidified ingot Grain clustering with strong internalorientation gradients
West and Adams, 1997
Zn alloy Hot dip galvanising Polycrystalline ‘grains’ with domains ofsame crystal orientation possiblyconnected in 3D
Semoroz et al., 2002a
Ni–18.7 at.-%Sn Containerless anomalouseutectic
Ni3Sn (near single crystalline):connected network; Ni: individualtextureless particles
Li et al., 2005b
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EBSD in an Al–1.25Mn alloy (Sun and Ryum, 1992)and in a beryllium alloy weld (Wright and Cotton,1995). A more or less strong ,100. texture was foundto gradually develop in X–750, CMSX–4 and CM186LCnickel base superalloys (Gandin et al., 1995; Ardakaniet al., 2000; D’Souza et al., 2000), in a twin roll cast Fe–3Si alloy (Takatani et al., 2000) and in austeniticstainless steels with simulated strip casting on Cusubstrates of controlled roughness (Hunter and Ferry,2002b and 2002c). In fact, the viscosity of the liquidcontrols the formation of secondary and tertiarydendrite arms (D’Souza et al., 2000). In case of lowviscosity, the competition between grains depends on therelative orientations of the primary dendrite arms [e.g. inCMSX–4 bicrystals, see (Wagner et al., 2004. Stanfordet al., 2004b)].
The origin of new undesirable ‘stray’ grains cannot beassessed from 2D measurements; in a laboratory nickelbase superalloy, serial sectioning coupled with EBSDmapping proved that stray grains nucleated at mouldwalls and not from dendrite fragmentation or constitu-tional undercooling (Stanford et al., 2004b). Stray grainformation was also characterised after electron beamand laser welding as a function of alloy composition,welding direction (with respect to crystal orientation),misorientation between welded parts and welding speed(Hirose et al., 2003; Vitek et al., 2003 and 2004).
A possibility to avoid small, undesirable grains is toimprove the shape of isothermal curves using appro-priate modelling coupled with EBSD experimentalvalidation (e.g. (Kermanpur et al., 2000) on IN738 alloyand (Carter et al., 2000) for innovative single crystalcasting design)
By allowing crystal orientation mapping over largeareas, EBSD is particularly well suited to studies of the
CET during solidification. The CET occurs due to localmodification of thermal and/or chemical or mechanicalconditions near the tips of growing dendrites. Examplesmay be found for eutectic solidification (see sectionon nucleation in eutectic solidification), and also fordendritic and cellular solidification (Table 6).
In hot dip galvanising, grains nucleate with no textureand competitive grain growth depends on both dendritegrowth directions and grain orientation with respect tothe solid/liquid interface (here the grain thickness ismuch lower than the in plane grain size) (Semoroz et al.,2002a and 2002b; Quiroga et al., 2004).
Such EBSD investigations have also been applied tothe more complex case of two phase solidification in Fe–19Cr–11Ni stainless steels, solidified by autogeneous gastungsten arc welding interrupted by liquid tin quenching(Inoue et al., 2000). Competitive grain growth in each ofthe ferrite and austenite phases was eventually governedby the local crystal orientation of the HAZ at the fusionline (Fig. 5).
Dendrite fragmentation and semisolid state
Dendrite fragmentation may appear during recalescenceor after partial melting of the seed in single crystal
4 Competitive grain growth during solidification, illu-
strated with EBSD grain boundary map of X–750Ni
base superalloy: after Gandin et al. (1995)
Table 6 Studies on CET using EBSD
Material Solidification process CET mechanismInformation provided byEBSD Ref.
CMSX–4 Ni basesuperalloy
Laser metal forming Constitutional undercooling Number of nuclei per unitvolume
Gaumannet al., 2001
IN718 Ni base superalloy Vacuum arc remelting Constitutional undercooling ‘Tree rings’ (solidificationdefects) consist of equiaxedgrains
Xu et al.,2002a
Ni–1B Electrostatic levitationmelting
Cell fragmentation(undercooling ,200 K)
Remaining solid is recognisedthanks to twin orientations
Li et al.,2005a
Ni–1B Electrostatic andelectromagnetic levitationmelting
Dendrite fragmentation(undercooling: 70–200 K)
Grain size and shape; rotationof fragments between non-fragmented dendrites
Li et al.,2006
Cu, Cu–1 at.-%Ag, Fe–10 at.-%Ni, low carbonsteel
Magnetic levitation melting Magnetic stirring: surfaceoscillation or bulk convectionin liquid sphere
Evidence of equiaxed areas Yasudaet al., 2005
5 Competitive grain growth during solidification of
duplex stainless steel in FA mode, after Inoue et al.
(2000). Unfavourably oriented c austenite grain 1 is
eventually replaced by grain 2. Unfavourably oriented
d ferrite grains (cases B, D, H) are eventually replaced
by d grains keeping (case B) or not keeping (cases D,
F, H) KS OR with growing c grain. As a result, lacy d
(resulting from KS OR, case A) is less frequently
found than vermicular d (no OR, cases C, E, G)
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74 International Materials Reviews 2007 VOL 52 NO 2
casting. For a high volume fraction of liquid, dendritefragments may rotate, so that a random texture is foundin that area after solidification [e.g. in CMSX–4 singlecrystals (Stanford et al., 2004a) and in Ni–1B lasermelted under electrostatic levitation (Li et al., 2005a)].In directionally solidified Ni–18.7 at.-%Sn having ananomalous eutectic microstructure, recalescence affectedthe Ni phase but not the Ni3Sn phase, so that Nifragments could not rotate easily and only LABs wereeventually observed with EBSD (Li et al., 2005b). Whenthe liquid fraction is rather low and no LABs areobserved between dendrite fragments, one may concludeto Ostwald ripening instead of coalescence [e.g. inreheated Al–7Si–0.3Mg alloy (Kliauga and Ferrante,2005)].
When a 6082 aluminium alloy is cast from thesemisolid state, grains tend to cluster into agglomerateshaving low energy (LAB or twin) boundaries. Thus,EBSD may distinguish ‘former’ aggregates after solidi-fication completion, showing intensive grain agglomera-tion if the mix is stirred, but very little in the absence ofstirring (Arnberg et al., 1999). Grain fragmentationduring agglomeration of silicon particles in thixoformedAl–30Si–5Cu–2Mg (wt-%) alloy has been shown usingEBSD to be of {111} type, leading to {111} agglomera-tion boundaries (Hogg and Atkinson, 2005). One mustnot forget, however, that EBSD does not distinguishbetween the former solid phase and the quenched liquid,which may solidify epitaxially on existing crystals, sothat great care must be taken while evaluating the grainboundary character distribution (Liu et al., 2005).
Solidification microtexture and resultingaverage textureSolidification microtexture
Orientation relationships between solid phases in assolidified microstructures have been investigated inmodified Al–Si alloys (see section on nucleation ineutectics). In as cast Al62.5–Cu25–Fe12.5 and Al65–Cu20–Fe15 alloys (Gui et al., 2001), no OR was found betweenl-Al13Fe4 single crystal dendrites and respectively theicosahedral quasicrystalline matrix and the polycrystal-line R-phase. In as welded austenitic stainless steel, theKurdjumov–Sachs OR (see ‘Appendix’) was foundbetween bcc d ferrite and fcc c austenite in lacy ferritemicrostructure (Inoue et al., 2000) whereas only acommon ,100. crystal direction, parallel to the growthdirection, was found between d and c phases in skeletalferrite microstructures (Inoue et al., 2000; Bouche et al.,2000). High resolution studies of solidification micro-textures are made possible thanks to EBSD. Forinstance, vacuum cast joints of white irons exhibitM7C3 carbides, surrounded by M3C with the samecrystal orientation, which are embedded in a ferritic ironmatrix (Wuhrer et al., 2004). Small, highly misorientedgrains have been observed in directionally solidified melttextured YBCO alloyed with YBa2CuO5 (Koblischka-Veneva et al., 2003; Koblischka et al., 2003). At (211)phase particles, the orientation of the (123) YBCOmatrix was strongly disturbed from its (001) texture,depending on the orientation of the individual (211)phases.
The EBSD is particularly useful if dendritic grains arevery coarse in size but with a very fine internal micro-structure. Few EBSD results addressed identification of
primary trunks from tertiary arms [e.g. in directionallysolidified Mg–9Al–1Zn alloy (Pettersen and Ryum,1989)]. Others focused on internal misorientation withindendritic grains. These may affect both morphology andinternal misorientations e.g. in ,112. dendrites of Al–4.3Cu–0.3Mg alloy in certain conditions (Henry et al.,1998b). In most cases, however, crystal orientationgradients are found within dendrites, due to fluid flowand shrinkage stresses (Doherty, 2003) or to gradient insolute concentration or to substrate induced thermalstresses (Semoroz et al., 2001). Some examples ofquantitative measurements using EBSD are given inTable 7.
Practical example: growth of ‘feathery’ grains
Very coarse, undesirable ‘feathery’ grains with fanshaped dendrites develop in certain conditions at theexpense of columnar grains in aluminium alloys. Theymostly appear during direct chill casting of alloyscontaining no refining elements, with high thermalgradients beyond the solidification front and underconvection (Rappaz and Henry, 1999). By EBSDmapping of wide areas at a fine scale compared to thedendrite arm width, Henry et al. showed that featherygrains were composed by parallel lamellae with alter-nating twin related crystal orientations; {111} bound-aries being alternatively (coherent and straight) and(incoherent and wavy) (Henry et al., 1997, 1998a and1998b). Within one main orientation, some subbound-aries were even found (Henry et al., 1997). Thecorresponding misorientation tended to accumulateleading to the fan shaped morphology (Henry et al.,1998b). The crystal direction parallel to the thermalgradient (,100.Al in columnar grains) was clearly,110.Al in feathery grains, with both ,110.Al
(possibly twinned) and ,100.Al secondary arms.,110.Al primary arms were twinned, the {111}twinning plane containing the growth direction.Coherent boundaries were found in arms parallel tothe thermal gradient. The detailed role of convection oncounterflow growth of ,110.Al twinned arms isreported elsewhere (Rappaz and Henry, 1999; Henryet al., 2004). Details about the anisotropy of primarydendrite arm spacing and 3D morphology can be foundin (Henry et al., 1998a).
Microtexture of eutectic and peritectic alloys
Eutectic and peritectic alloys often exhibit coarsecolonies of fine lamellae, so that high spatial resolutionis required over wide areas. Sample preparation is madedifficult by the difference in electrochemical and hard-ness properties of the various phases, so that EBSDstudies are still scarce. Some results on ORs found ineutectic and peritectic materials are illustrated inTable 8.
Average texture of as solidified products
The average texture is in many cases preferablyinvestigated by XRD. However, when the grain sizeexceeds a certain value (typically 50–100 mm), XRD canno longer provide statistically significant data. Neutrondiffraction can be used for grains up to y1 mm in size.For even coarser grains and for heterogeneous materialssuch as welded joints, EBSD can still be used thanks tostage motion in the SEM. In EBSD, the orientationdistribution function (ODF) is then directly calculated
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Table 7 Electron backscatter diffraction measurements of internal misorientation in dendritic grains
Material Solidification processMisorientations measuredwith EBSD Ref.
Al–0.15Fe–0.07Si Direct chill casting Many misorientations upto 5u within dendrites
Samajdar and Doherty,1994
CMSX–4 Ni base single crystalsuperalloy
CO2 laser metal forming After 5 passes, arms becometrunks with 5–10u misorientationbetween dendrites
Gaumann et al., 1999
CMSX–4 Ni base single crystalsuperalloy
CO2 laser epitaxial cladding Slight misorientations betweendendrite trunks are viewedthanks to specifically developeddata processing
Cleton et al., 1999
CMSX–4 and CMSX–10N Nibase single crystal superalloys
Directional casting of turbineblades
Steady state growth: internalmisorientations of 2–3u do notcumulate; non-steady stategrowth: successive branchingcumulates misorientation up to y6u
D’Souza et al., 2005;Newell et al., 2005
99.9% pure Al Ingot casting (columnar zone) Misorientation of 4–10u over 3 mmalong the growth direction; internalLABs of 6–8u
Bhattacharyya et al.,2001
Al–Zn–Si alloy Hot dip galvanising Dendrite arms misoriented by4–10u mm21. Misorientations up to5u within grains, cumulate up to35u in coarser grains
Semoroz et al., 2001
High Cr white cast irons Electromagnetic stirring andcasting
Austenite grains are clusters ofsubgrains and internal LABs aresolute enriched and could originatefrom dendrite arm bending
Yang et al., 2003
Zn–0.2Al–0.15Sb Hot dip galvanising One nucleus only for dendrites in agiven grain; orientation domainsrelated by simple ORs in‘polycrystalline’ grains
Semoroz et al., 2002a
Table 8 Electron backscatter diffraction measurements of orientation relationships in as solidified eutectic andperitectic microstructures*
MaterialSolidificationmode Microstructure OR between phases Growth direction Ref.
Ni–Cr–(Si) whitecast iron
Slow cooling(0.017 K s–1)
E for 0%Si, P for2%Si
0%Si: no OR betweenM3C formed at a givenM7C3
Strong [0001]M7C3
for 2%Si; no texturein M3C for 0%Si
Randle andLaird, II 1993
White iron Strip casting E Plate like Fe3C: [001]Fe3C//
normal direction // [001]a
[001]Fe3C// plane of
platesSong et al.,2003a
V–13 at.-%Si Directional E (broken lamellar) (011)V3Si // (112)V // lamellaeboundaries
[100]V3Si // [111]V Bei et al.,2004
Cr–16 at.-%Si Directional E (lamellar) (011)Cr3Si // (123)Cr // lamellaeboundaries, (001)Cr3Si // (011)Cr //lamellae boundaries
[100]Cr3Si // [111]Cr (closepacked directions)
Bei et al.,2003
Ni–45.5Al–9Mo(at.-%)
Directional E (NiAl matrix z
Mo fibres){011}Mo // {011}NiAl // interfaces ,100.Mo // ,100.NiAl Bei and
George, 2005Mg–33Al–Sr(wt-%)
Directional E (lamellar) (1101)Mg // (101)Mg17Al1210u from
interface[1120]Mg // [111]Mg17Al12
2–15u from growthdirection
Guldberg andRyum, 2000
Nb–22 at.-%Si Arc melting inCu crucible
E {111}Nb3Si // {111}Nb (one areastudied)
Drawin et al.,2005
Nb–16Si–1.5Zr(at.-%)
Arc melting inCu crucible
E No OR between Nb and Nb3Si Miura et al.,2005c
Al–7Si–(0–150 ppmSr)
Ingot casting E No Sr: two twin related Si variants;with Sr: various twin related Sivariants with one common ,110.
per colony, often parallel to ,100.
or ,110.Al; no clear OR betweenAl and Si
Heiberg andArnberg, 2001
Al2O3–19 mol.-%YAG
Directional(fibres)
E ,0001.Al2O3// ,112.YAG Nakai et al.,
2005Al2O3–16%YAG–18 mol.-%ZrO2
Micropulling E (lamellar) Three intricate single crystals pereutectic grain, ,2110.Al2O3
//,100.ZrO2
, YAG is much lesstextured
Murayamaet al., 2004a
Al2O3– 16%YAG–19 mol.-%ZrO2
Directional(rods)
E (‘geometric’ or‘Chinese scripts’)
No clear OR between fibretextured Al2O3 and ZrO2
Murayamaet al., 2004b
*E: eutectic; P: peritectic; YAG: yttrium–aluminium garnet.
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76 International Materials Reviews 2007 VOL 52 NO 2
from individual orientations by taking one point pergrain or by taking all data points and weightingorientation data with the respective area fraction ofgrains. If the grain size is low enough, EBSD and XRDresults can be successfully compared (e.g. Slamova et al.,2003; Saito et al., 2004). One must, however, keep inmind that EBSD is a surface analysis technique.Sectioning effects may be significant when grains arenot equiaxed, which is frequently the case in solidifica-tion microstructures.
Spatial information given by EBSD is necessary whenthe heat flow direction is not uniform within the part,e.g. in thin strip casting (Table 9) and in welded joints(Table 10). In some instances, EBSD is used as a com-plementary method to evaluate GBs of samples whilethe average texture is determined by XRD (Park et al.,1999). Some examples obtained in as solidified bulkmaterials are listed in Table 11.
A strong average texture induces particular ORsbetween grains, as e.g. in a ,100. textured directionally
Table 9 Solidification textures determined with EBSD in thin strip castings and hot dip coatings
Material Microstructure Average texture Ref.
Cast iron Grey, compacted ornodular
Inoculated materials: notexture in austenite andthus no texture in resulting ferrite
Campos et al., 2005
0.1%C steel (dendritic growth) Strong {100} texture Umezawa et al., 2003Fe–3Si ‘Upper’ and ‘lower’
columnar zones versusmid-thickness
,100. // local thermalgradient versus randomly textured
Takatani et al., 2000
Fe–15 at.-%Ga Strong ,100. texture withmany LABs
Saito et al., 2004
AISI 409 ferritic stainless steel Columnar zone ,001. fibre (ferrite) Ferry and Hunter, 2002AISI 304 austenitic stainless steel Columnar zone ,001. fibre (austenite) Ferry and Hunter, 2002AISI 409 ferritic stainless steel(simulated thin strip casting)
Chill zone and growthzone
Ti free (respectively Ti containing)steel: random texture (respectively,001. fibre) in chill zone, ,001.
increases in growth zone in bothcases
Hunter and Ferry, 2002a
AISI 304 austenitic stainless steel(simulated thin strip casting)
Chill zone and growthzone
,001. fibre increases duringgrowth
Hunter and Ferry, 2002b
Sn–3.5Ag and other Pb free solderalloys
Multicrystalline solderjoints
,110]Sn // heat flow, mainly‘special’ GBs
Telang et al., 2002;Telang and Bieler, 2005a,2005c
Al–Mg 5052 and 5182 alloys Elongated and equiaxedgrains
Less cube component than indirect chill castings
Slamova et al., 2003
Al–43.4Zn–1.6Si (hot dip coating) No average texture: the coatingdoes not inherit that of steelsubstrate
Semoroz et al., 2001
Zn–0.2Al–0.15Sb (hot dip coating) 43% of grains have their {0001}planes less than 22.5u from thefree surface due to constrainedgrowth
Semoroz et al., 2002a
Table 10 Average solidification textures determined with EBSD after welding or laser metal forming*
Material Process Microstructure Average texture Ref.
5182 and 6111 Al alloys Laser welding Columnar andequiaxed zones
,100. for columnar,random for equiaxed
Hector et al., 2004
Austenitic stainless steel GTA welding then liquid tinquenching
Cellular ,001. (ferrite) Inoue et al., 1995and 1998
AISI 409 ferritic stainlesssteel with Monel Ni–30Cu–Mn
Dissimilar GTA welded joints Growth area ,100. (austenite) Nelson et al.,1999b
Ni alloy 182 Multipass MMA welding Columnar ,100. Scott et al., 2005Ni alloy 82 Multipass GTA welding with
or without magnetic stirringStraight solidificationfront under magneticstirring, curved frontotherwise
,100. under magneticfield, random otherwise
Kokawa et al.,1999
Mg and AZ31, AZ61 andAZ91 Mg alloys
Autogeneous EB welding Weak texture Su et al., 2002
Mg alloys Autogeneous GTA welding Weak texture Wu et al., 2004A286 Fe based superalloy CO2 laser welding Weld metal centre Rotated cube Weiß et al., 2002CMSX–4 Ni base superalloy CO2 laser welding Welding at 30u from
,001.: many smallgrains
,001. // local thermalgradient
Hirose et al., 2003
CMSX–4 Ni base superalloy Multipass CO2 laser metalforming
Substrate and deposit ,001. fibre within 5u forsubstrate and within 10ufor deposit
Gaumann et al.,1999
*GTA: gas tungsten arc; MMA: manual metal arc; EB: electron beam.
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International Materials Reviews 2007 VOL 52 NO 2 77
solidified nickel base alloy ingot (West and Adams,1997); many subboundaries developed and grainsclustered into coarse entities having internal LABs andstrong orientation gradients. The size of such entitiesgrew faster than the grain size itself, due to stronglytextured (competitive) grain growth. In directionallysolidified CMSX–4 and CM186LC nickel base super-alloys, misorientation axes got close to the ,100. fibreduring growth; many LABs were also found, but‘special’ GBs such as S5, S13a and S17a (all misor-ientation around ,100. axes) were not favoured(Ardakani et al., 2000).
Hot cracking and GB resistanceHot cracking during solidification is of major practicalconcern. The EBSD was mainly applied to problemsassociated with welding. In nickel base alloy 718, HAZliquation was shown to occur at random GBs but not atcoincidence site lattice (CSL) boundaries, according tothe low sensitivity of CSL boundaries to boron atomsegregation (Guo et al., 1998) and to boride and carbideprecipitation (Qian and Lippold, 2003). In the [0001]Cu6Sn5 layer of solder interconnects, many HABsapparently provided penetration channels for the solderalloy (Lee et al., 2001). In free machining laser or gas
tungsten arc welded stainless steels, only alloys prone tofirst solidify into austenite were prone to hot cracking,according to the higher sensitivity of austenite toboundary segregation and cracking than that of ferrite(Brooks et al., 2003). In gadolinium containing Ni basealloy welded by electron beam or gas tungsten arcprocesses, epitaxial solidification evidenced with EBSDprovided a good resistance to hot cracking according toVarestraint tests (Robino et al., 2003).
Welding of nickel base single crystal superalloys maylead to hot cracking if the single crystals to be joined aremisaligned from each other or from the weldingdirection. The EBSD helped to establish adequatewelding conditions and to identify, if any, the highlymisoriented, newly nucleated grains that can induce hotcracking. Examples are found in laser welded CMSX–4(Hirose et al., 2003), in laser and electron beam weldedRene N5 (Vitek et al., 2003 and 2004) and in laserwelded MC2 (Wang et al., 2004) single crystals orbicrystals. Here, EBSD data on base metal orientationand WM microtexture were combined with analyticalmodels involving, e.g. local undercooling at grainboundaries.
Hot cracking in Al–4Cu alloy was also investigated byEBSD in terms of wetting of GBs by the remaining
Table 11 Solidification textures determined with EBSD in bulk materials (results of Grant et al. (1986) and Hanada et al.(1986) were obtained with pioneering electron channelling pattern technique in SEM)
Material Process Microstructure Average texture Ref.
Cu Directional solidification Columnar ,001. Grant et al., 1986AA1100 Al alloy Ingot casting Almost equiaxed Random Wang et al., 1990Al–1.25Mn Directional solidification Cellular ,001. Sun and Ryum, 1992AA 5052 and 5182 Al–Mgalloys
Direct chill and twin rollcasting
Equiaxed Copper z Brass z
cube z other texturecomponents
Slamova et al., 2003
Ti–6Al–4V Vacuum arc remelted ingotcasting
Columnar ,100. (b phase) Glavicic et al., 2003c
Ti–48Al–2Cr–2Nb (at.-%) Ingot casting Columnar ,0001. of a phase Dupont et al., 1996Ni3Al Induction melting Columnar ,001. Hanada et al., 1986X–750 Ni base superalloy Directional solidification Columnar ,100. increases during
growthGandin et al., 1995
Fe–18.4 at.-%Ga Zone melting Columnar ,100. (ferrite) Summers et al., 2004CMSX–4 Ni base superalloy Directional solidification Columnar ,100. fibre Carter et al., 2000Duplex stainless steel Statically cast Columnar ,100. (ferrite) Calonne et al., 2000AISI 430 ferritic stainless steel Continuously cast slab Columnar zone
equiaxed zone{001},uv0. (columnar),random (equiaxed)
Hamada et al., 2003
Nb–33Ti–16Si (at.-%) Czochralski directionalsolidification
Cellular Weak ,113. for (Nb);,001. (almost singlecrystalline) for (Nb,Ti)3Si
Sutliff and Bewlay,1996
Nb–16 at.-%Si Arc melted ingot casting (Nb) dendritesz (Nb)/Nb3Sieutectic
[001] for eutectic Nb3Si;(Nb) randomly textured
Drawin et al., 2005
Nb–22 at.-%Si Arc melted ingot casting Nb3Si dendritesz (Nb)/Nb3Sieutectic
[001] for dendritic Nb3Si;(Nb) textured due to ORwith Nb3Si
Drawin et al., 2005
Cu6Sn5 Soldering Cu6Sn5 scalloplike grains
[0001] Lee et al., 2001
ZrB2/ZrC Spark plasma sintering ZrB2 weakly textured, ZrCrandomly textured
Shim et al., 2002
Al/Al3Ti Centrifugal solid particlecasting
Coarse Al grainsand Al3Ti platelets
Both Al and Al3Ti arerandomly textured
Watanabe et al.,2002
V–V3Si Directional solidification Eutectic ,111.V//,100.V3Si Bei et al., 2004Al2O3–YAG–ZrO2 eutectic Modified micropulling ‘Chinese scripts’
‘geometric’,2110. or ,1010.Al2O3
,,001. or ,220.ZrO2
,possibly ,100. or,111.YAG
,0001. or ,1010.Al2O3,
,001. or ,220.ZrO2,
,100. or ,111.YAG
Murayama et al.,2004b; Lee et al.,2005
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78 International Materials Reviews 2007 VOL 52 NO 2
liquid as a function of barium content and GB misor-ientation. By decreasing the liquid/solid interfacialenergy (deduced from minimum misorientation angleof wetted GBs and Read–Shockley equations) and thusdelaying grain coalescence, barium strongly increasedhot cracking sensitivity of this alloy (Fallet et al., 2006).
Electron backscatter diffraction andsolidification: summaryElectron backscatter diffraction is of particular interestin solidification investigations because of its ability tocharacterise the microstructures at multiple scales withsuitable statistical significance, from (sometimes very)coarse grains down to submicrometre sized phases suchas nucleating agents and eutectic colonies. Its angularand spatial resolution is generally high enough for thatpurpose. Geometric and statistical effects due tosectioning and sampling effects must be taken intoaccount or, if possible, avoided by carrying out 3Dinvestigations (see ‘Discussion’).
Solid state phase transformations
Orientation relationships between phasesMatrix phase transformations
The EBSD was generally used in this field to confirmTEM or XRD data or, alternatively, to get statisticallyreliable results. A huge variety of materials and phasetransformations have been addressed with EBSD asillustrated in Table 12 for ceramics, in Table 13 for non-ferrous metals and in Table 14 for intermetallic phasesand quasicrystals. Details about the most frequentlyobserved ORs are reported in the Appendix.
Most ORs were investigated at the scale of individualphases or by comparing the pole figures of productphases (formed from a single grain of the parent phase)to pole figures calculated using known ORs. Theysuccessfully compare with TEM data. Only few inves-tigations used the average pole figures (Boehlert andBingert, 2001).
Like eutectic microstructures, products of eutectoiddecomposition usually exhibit fine scale features and inmany cases only a few colonies may be observed in asingle TEM thin foil, whereas large areas may bescanned with EBSD. Four examples of results obtainedusing EBSD are given here:
(i) in Nb–(16–25)Si¡1.5Zr (at.-%), the Nb3Si partof the (NbzNb3Si) eutectic further decomposesupon cooling into tetragonal a-Nb5Si3zbcc(Nb). In binary Nb–Si alloys, eutectoid (Nb) issingle crystalline within a given eutectoid colonyand may share the same crystal orientation asthat of neighbouring eutectic (Nb) (Drawinet al., 2005). No OR was found between (Nb)
and a phases for 16 at.-%Si alloy, while in Nb–22 at.-%Si the authors found (111)Nb//(100)a
and (011)Nb//(011)a or (001)a (Drawin et al.,2005). In ternary alloys, occurrence and natureof the OR depend on chemical composition andheat treatment (Miura et al., 2005b and 2005c)
(ii) in Al–36Mo–17Ti (at.-%), the b (A2) phasetransforms into Al3Ti (DO22) and Mo3Al (A15)phases. An OR was found between b and DO22,and between b and A15 phases (Miura et al.,2005a). Such OR disappeared after hot defor-mation (Miura et al., 2005a) probably due to amemory effect on variants due to the orderedstructure of both product phases (Miura et al.,2004)
(iii) in Fe–12Mn–0.8C¡0.3V steels, decompositionof fcc austenite c into bcc ferrite a andtetragonal Fe3C cementite depends on thecomplex sequence of nucleation events (aformed at VC carbides formed at MnSsulphides) so that the KS OR between c anda is not systematically observed (Guo et al.,2001 and 2002)
(iv) in Mn–Si bearing hypo- and hypereutectoidsteels decomposed into pearlite under smallundercooling, the OR between a and Fe3Cstrongly depends on the 3D topology of phases.Given a colony nucleated at a c GB, the Pitsch–Petch OR (see ‘Appendix’) prevailed if the aphase of the pearlite colony was totally isolatedfrom the c grain where the colony did notdevelop, while the Bagaryatsky OR (see‘Appendix’) prevailed otherwise (Fig. 6). Inboth cases, Fe3C crystals at c GBs and in thecolony were found to share the same crystalorientation (Mangan and Shiflet, 1999).
In the last two examples, only association betweencareful serial sectioning and EBSD analysis allowed toget the key results. This point will be further addressedin the ‘Discussion’ section.
Precipitation
Most ORs between matrix and precipitates wereinvestigated with TEM. However, if precipitates arescarcely distributed (e.g. at boundaries of coarse grains)or if no high angular accuracy is required, EBSD mayreadily provide relevant data (Table 15). The EBSD isalso useful to study interactions between precipitationand damage cavitation, as much less sample preparationartefacts are usually induced than by TEM thin foilpreparation. Discontinuous precipitation has also beeninvestigated with EBSD. Results illustrated in Table 16are generally complementary to data obtained withTEM, although the evolution of the OR along the
Table 12 Orientation relationships determined with EBSD in solid state phase transformations of ceramic materials
Material Parent phase Product phase Orientation relationships Ref.
(001) MgO coatedwith In2O3
MgOzIn2O3 MgIn2O4 [001]MgO//[111]MgIn2O4and
[110]MgO//[110]MgIn2O4; no
OR between In2O3 and MgIn2O4
Korte et al., 2000
LaNbO4 Tetragonal Monoclinic Misorientation of 94u around[010] between product phasedomains
Jian and Wayman, 1995
Iron oxides Hematite Fe2O3 (H) Magnetite Fe3O4 (M) ,0001.H//,111.M and,1010.H//,110.M
Piazolo et al., 2004b
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 79
growing colonies might be analysed with EBSD only,especially over large distances (e.g. the last item inTable 16).
Decomposition of ferrite and austenite in ferrous alloys
Austenite (fcc c) and ferrite (bcc a or d) phases have beenwidely studied with EBSD. Several ‘classical’ ORsclustered around the 45u,001. Bain OR are recalledin the ‘Appendix’. There is still much debate about therelationship between OR and the displacive or diffusivecharacter of phase transformations in steels, in parti-cular concerning acicular ferrite, Widmanstatten ferrite
sideplates and bainite. This debate is well beyond thescope of the present review: to illustrate the variety andinterest of EBSD data about these as yet incompletelyunderstood phase transformations.
In duplex stainless steels, both d and c phases are stillpresent at room temperature, which facilitates the studyof their OR. As solidified microstructures are verycoarse, so that TEM based methods are not well suited.Note that c forms at high temperature (.1000uC) andthat high thermal stresses are expected to arise fromcooling due to the difference in thermal expansioncoefficient between c and d. The effect of such stresses on
Table 13 Orientation relationships determined with EBSD in solid state phase transformations of non-ferrous and Ni–Femetal alloys
Material Parent phaseProductphase Analysis scale Orientation relationships Ref.
Ti sheet a (hcp) b (bcc) Grains and GBs (in situ) Burgers; {334}b midribparallel to the commonclose packed plane
Seward et al.,2004
Ti–6Al–4V b (bcc) a (hcp) Macrozones Burgers for both primaryand secondary a
Le Biavantet al., 2002
Ti–8Al–(0–20)V b (bcc) a (hcp) Along one GB with gradientof chemical composition
Burgers (with oneexception)
Banerjeeet al., 2004
Ti–6Al–2Sn–4Zr–6Mo b (bcc) a (hcp) a colonies near b GBs Burgers Bhattacharyyaet al., 2003
Ti–5Ta–1.8Nb b (bcc) a (hcp) Widmanstatten colonies atGBs and primary a
More or less close toBurgers, neighbouringWidmanstatten and primarya phases having samecrystal orientation
Karthikeyanet al., 2005
Near-a a IMI834 Tialloy
b (bcc) a (hcp) Colonies of a phase Loss of Burgers OR;common close packedplanes preserved afterspheroidisation of the aphase
Germainet al., 2005b
Co b (fcc) a (hcp) Individual phases Near NW Wright et al.,2005
Cu–42 wt-%Zn b (bcc) a (fcc) At least b grains KS (from unpublishedpole figures)
Sakata et al.,2000
Cu–40Zn b (bcc) a (fcc) Widmanstatten colonies 1.7u from KS; no strict OR Stanford andBate, 2005
Cu–11 wt-%Ag b (Cu) a (Ag) Individual phases Cube–cube, i.e. samecrystal orientation
Li et al.,1994
Cu–Zn–Al b (bcc) a (fcc) Individual variants ofbainitic a phase
Pitsch Marukawaet al., 2000
Pu–2 at.-%Ga e (bcc) d (fcc) A few d variants One common ,110.d:possibly KS or NW
Boehlertet al., 2003b
Cu–12.55Al–4.84Ni(wt-%) shape memoryalloy (SMA)
P (DO3) 2H(orthorhombic)
Microtexture of diamondshaped entities
{001}2H // {110}P and[010]2H // ,001.P
Chen et al.,2000
Cu–7.3Al–8.5Mn (wt-%)SMA
L21 18R1
(monoclinic)Microtexture of lensshaped entities
Twin misorientationrelationships (MRs)between product variants
Wang et al.,2002
Fe–27.5Ni–17.7Co–3.8TiSMA
c (fcc) a9 (bcc) Local scale before andafter phase transformation
80% of ORs at (5ufrom NW
Bruckneret al., 1999
Fe–32 at.-%Ni bicrystals c (fcc) a9 (bcc) Near the GB NW Ueda et al.,2001a
Fe–32.85 wt-%Ni c (fcc) a9 (bcc) Individual variants Near GT close to themidrib; near KS close toretained austenite;gradient of OR in between
Shibataet al., 2005
Fe–29.6 wt-%Ni c (fcc) a9 (bcc) Former austenite grains,individual variants
Near NW Kitaharaet al., 2005
Ni/Ti/Ni multilayers c-Ni (fcc), a-Ti(hcp)
NiTi (B2) Average texture Probably KS for c/B2 andprobably Burgers for a/B2
Inoue et al.,2003
Ti–5Al–2Sn–4Zr–4Mo–2Cr–1Fe (b-Cez) alloy
b (bcc) a0
(orthorhombic)Individual crystals (twoperpendicular sections)
Habit plane 2u in averagefrom [312]
Zimmermannand Humbert,2002
Ni–(36–38)Al (at.-%) b (B2) indexedas bcc
c9 (L12)indexed as fcc
Individual crystals KS Sakata et al.,2001
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80 International Materials Reviews 2007 VOL 52 NO 2
the OR (which is measured at room temperature) is notknown yet. The OR was reported to be close to KS in ascast and as welded structures (Inoue et al., 1998; Pinol-Juez et al., 2000; Gourgues et al., 2004). In someinstances a unique OR, yet different from the wellestablished ones, was identified (Nolze, 2004 and 2006).Widmanstatten austenite tends to grow along thecommon close packed direction of the KS OR and toshare the same orientation as neighbouring allotrio-morphic (primary) austenite phase (Inoue et al., 1998),possibly with only one adjacent d grain into which it isgrowing (Pinol-Juez et al., 2000). Even massive austeniteforms in near KS OR with the parent d grain, at leastin autogeneous gas tungsten arc welded, then liquidtin quenched 21Cr–9.3Ni stainless steel (Inoue et al.,1995). Gas nitrided 22.5Cr–5.4Ni–1.9Mn–3.0Mo–0.16N
stainless steel, however, undergoes subsurface growth ofneedle shaped austenite having close packed planes aty15u from those of d ferrite (however, ferrite andaustenite textures were not measured in the samesample) (Tschiptschin et al., 2002).
Decomposition of ferrite into austenite has also beenstudied during heating in ferritic steels. To keep someaustenite phase retained after cooling at the end ofexperiments, the low carbon steel studied (0.06C) waswrapped into either nickel or austenitic stainless steelfoils before heating up to y900uC for a few minutes, tostabilise high temperature c by element diffusion(Bruckner et al., 2001; Park et al., 2002). The KS ORwas not strictly followed (departing by up to 10 or even25u). However, artefacts due to surface effects and to thelocal change in chemical composition still need further
Table 14 Orientation relationships determined with EBSD in solid state phase transformations of intermetallic alloysand quasicrystals
Material Phase transformation Analysis scale Orientation relationships Ref.
Al62.5–x–Cu25.3–Fe12.2–Bx Secondary phases in facecentred icosahedral (fci)matrix during sintering
Individual phases No OR between fci andsecondary phases
Brien et al.,2004
Ni base superalloy coatedwith NiAl
Pack cementation Individual phases(L12 phase indexedas fcc structure andB2 phase indexedas bcc structure)
{111}L12//{110}B2 and in
that plane either,112.L12
//,110.B2 or,011.L12
//,111.B2
Wollmer et al.,2003
Ni3Al coated with NiAl Pack cementation Individual phases (001)[110]L12//either
(011),111.B2 or(101),111.B2 or(014),041.B2
Zaefferer andGlatzel, 2002
Ni3Al coated with NiAl Precipitation of body centredtetragonal (bct) W2Ni
Individual phases (110)bct//(111)L12and
[001]bct//[110]L12
Zaefferer andGlatzel, 2002
Ti–25Al–24Nb (at.-%) a2 into O A few O variants Twin MR betweenvariants consistentwith a2 to O phasetransformation
Li et al., 2004
Ti–48Al–2Cr–2Nb (at.-%) a into massive c (indexedas fcc)
Individual phases (0001)a2//{111}c and
,1120.a2//,110.c z
identification of orderdomains
Pouchou et al.,2004a and2004b
Ti–48Al–2Cr–2Nb (at.-%) a (indexed as hcp) into c(indexed as fcc)
Individual phases Burgers Dupont et al.,1996
Ni–45 at.-%Al a2 into cz spheroidisation ofa2
Individual phases (0001)a2//{111}c and
maybe ,1120.a2//,110.c
(c indexed as fcc structure)
Buque andAppel 2002
Ti–25Al–10Nb–2V–1Mo (at.-%) b (B2) into a92 (DO19) Individual phases (011)b//(0001)a92and
[111]b//[2110]a92
Yang et al.,2003
Ti–46.5 at.-%Al hcp a into massive L10 czDO19 a2 (at GBs)
Lamellar grains Burgers OR with onegrain, growth into anotherone
Wang et al.,2002
Ni–36 at.-%Al b into c9 (partial, at GBs) Individual phases ,5u from KS with oneb grain, no OR with theother neighbouring b grain
Sakata et al.,2001
(Fe,Al,Ni)–10Cr a into A2 z B2 (ordering) Individual phases Cube–cube OR, raftingalong ,100. direction
Stallybrassand Sauthoff,2004
Ti–46.8Al–1.7Cr–1.8Nb (at.-%) a into c (L10) z a2 (DO19) Individual phases Blackburn (see ‘Appendix’) Dey et al.,2005
Udimet 720LI Ni basesuperalloy
Interactions betweenrecrystallisation of c matrixand c9 precipitates
Former c grains Loss of OR between c andc9: c9 keeps the cube–cubeOR with parent c grain orturns to twin OR with therecrystallised c grain
Lindsley andPierron, 2000
Mo z (210) single crystalMoSi2
Diffusion couple, formationof Mo5Si3
Individual phases [001] textured Mo5Si3(cross-sectional observations)
Tortorici andDayananda,1999
Ti–26Al–27Nb–0.03O bcc into O Average texture Results consistent with thefollowing OR: (001)O//(110)bcc
and [110]O//[111]bcc
Boehlert andBingert, 2001
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International Materials Reviews 2007 VOL 52 NO 2 81
analysis. Synthetic Fe–4Cr–0.008C heated up to 820uCfor 20 h and mapped with EBSD both before and afterthe heat treatment showed, in view of the interfacemigration velocity, that starting a phase and newly
formed c phase shared at least one common closepacked plane (Watanabe et al., 2005b).
The transformation of austenite into proeutectoidFe3C cementite was investigated with EBSD togetherwith 3D considerations in high-purity Fe–1.34C–13.1Mn (Mangan et al., 1999; Kral and Spanos, 2003).Seventy five per cent of the ‘dendritic’ Fe3C particlesexhibited the Pitsch, Farooque–Edmonds orThompson–Howell ORs with the neighbouring c matrix.Monolithic Fe3C plates also exhibited the Pitsch OR,while conglomerates of parallel laths exhibited theFarooque–Edmonds OR. Many cementite crystals stillexhibited no OR with neighbouring austenite. This wasattributed to nucleation outside the investigated volume.Similar results were found with Widmanstatten cemen-tite plates grown from austenite in an Fe–0.8C–12.3Mnsteel (Mangan et al., 1997).
Decomposition of austenite into ferrite, bainite ormartensite is a key point in the processing of most highstrength steel grades. Therefore, many results nowcomplete the database acquired with TEM basedmethods for many years. One major difficulty is theabsence of retained austenite in many steels of practicalinterest. Details are given here for three kinds of phasetransformations according to their temperature range:
(i) martensite
(ii) ‘high temperature’ phase transformations, i.e.formation of idiomorphic, allotriomorphic andWidmanstatten ferrite
(a) (b) (c)
6 Interrelationships between morphology and ferrite–
cementite OR in pearlite. Pitsch–Petch OR in a
hypoeutectoid Fe–C and b hypereutectoid Fe–C if
pearlitic ferrite is not connected to austenite grain into
which it is not growing; c Bagaryatsky OR if pearlitic
ferrite appears disconnected from neighbouring auste-
nite grain (top) but is in fact connected as shown by
further serial sectioning (bottom). After Mangan and
Shiflet (1999)
Table 15 Determination of OR between matrix and precipitates using EBSD
Material Matrix Precipitate Results Ref.
Meteoritic minerals a (bcc) (Fe,Ni)3P No simple OR (consistentlywith TEM results)
Geist et al., 2005
Sintered Ti–4Fe–7.3MozTiB b (bcc) TiB (Widmanstatten) [010] growth of needles (OR:by TEM)
Feng et al., 2005
Zr cladded Zircaloy–2 in variousmetallurgical states
a (hcp) d (fcc) hydrides In general (0001)a//{111}d and,1120.a//,110.d; GB d with suchOR with only one neighbouringa grain; little effect of residual stresses
Une et al., 2004;Une and Ishimoto, 2006
SAF2507 superduplex stainlesssteel
c (fcc) s (decompositionof bcc d)
24u from the Nenno OR: here [101]c//[310]s Dobransky et al., 2004
Super austenitic stainless steel c (fcc) s (111)c//(110)s and [211]c//[110]s (notthe Nenno OR)
Lewis et al., 2006
Austenitic stainless steel c (fcc) M23C6 Cube–cube OR with one c grain,growth with no OR into the neighbouringc grain
Hong et al., 2001
Austenitic stainless steel c (fcc) M23C6 50% of particles in cube–cube OR,,20% next to creep cavities
Abdul Wahab and Kral,2005
Table 16 Determination of OR using EBSD for discontinuous precipitation
Material Parent phase Product phases OR Ref.
Cu–11 wt-%Ag b(Cu) (fcc) b(Cu)za(Ag) (all fcc) Cube–cube Li et al., 1994Cu–4 wt-%Ti a (fcc) a(Cu)zCu4Ti (orthorhombic) (111)a//(010)O and
[011]a//[501]O or(511)a//(010)O and[011]a//[501]O
Mangan and Shiflet, 1997
Re and Ru rich nickelbase alloy single crystal
czc9 (samecrystal orientation)
czc9ztopologicallyclose-packed phases
New c and c9 phaseskeep cube–cube OR
Lavigne et al., 2001 and2004
Nitrided Fe–(1–3)Cr a (bcc) azCrN a first cube–cube withparent a grain but twinboundaries developduring growth becauseof high volume increaseduring phase transformation
Sennour et al., 2004
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82 International Materials Reviews 2007 VOL 52 NO 2
(iii) ‘intermediate temperature’ phase transforma-tions, i.e. acicular and bainitic ferrite.
Martensite
Martensite results from a displacive phase transforma-tion and develops with a given OR with the parentaustenite grain into which it grows. Austenite formingelements may be added to steel composition to retainsome austenite, to facilitate investigation of the OR.Otherwise, the MRs between variants formed within agiven former austenite grain may be compared to thosecalculated using known ORs, e.g. KS/KS or NW/NWMRs. The most frequently (tried and) found ORs wereKS and NW for a9 (near bcc) martensite and Burgers forhcp e martensite (Table 17).
It is then possible to derive the crystal orientation ofparent austenite from those of martensite variantsformed within that austenite grain. This can beperformed by either graphical methods using {100} polefigures (Gourgues et al., 2000; Lambert-Perlade et al.,2004a; Cabus et al., 2004b) (Fig. 7) or by algebraiccalculations from MRs between variants (Suh et al.,2002). The former method makes no particular assump-tion about the OR (except that it is close to the BainOR), while the latter method assumes that the KS OR isstrictly followed, which is never exactly the case.
Allotriomorphic, idiomorphic and Widmanstatten ferrite
Orientation relationships between fcc and bcc phaseshave been investigated with high resolution EBSD inplessite, a mixture of fcc (taenite) and bcc (kamacite)phases found in iron and nickel rich meteorites. The ORvaries between KS and NW according to the particularmicrostructure and chemical composition of phases(Nolze and Geist, 2004). As kamacite may be deformed,such determination may be difficult; close packed planesare not strictly parallel and continuous variations of thelocal OR may be found (Nolze et al., 2005). In Ni–43 wt-%Cr alloy solution (of course not a ferrous alloy)annealed at 1200uC for 1 h then aged at 1000uC, bcc Crclusters exhibited close packed planes (respectively closepacked directions) 1.27u (respectively 2u) in average fromthose of the fcc (Ni) matrix, i.e. the OR was up to 6u offKS and NW ORs (Adachi et al., 2005), irrespective ofthe deformation applied before ageing (Adachi andTsuzaki, 2005).
The OR between primary ferrite and austenite wasgenerally determined by deriving the orientation ofaustenite from those of several neighbouring martensitevariants after interrupted heat treatments (Suh et al.,2002). The assumption of KS OR is, however, asimplification as long as up to 10–20% (respectively50%) of the MRs between martensite variants departed
from the KS/KS one by more that 10u (respectively 5u)(Suh et al., 2002; Cho et al., 2002a; Hernandez et al.,2003). Thus, the method cannot be highly accurate.Even a shift of only a few degrees off the strict KS (orNW) OR leads to an uncertainty double as high forMRs between product phases which can reach at least 8u(Gardiola et al., 2003). In fact, near KS or NW ORswere generally found between primary ferrite grains and(at least) one of the surrounding parent austenite grains(Suh et al., 2002; Gardiola et al., 2003). Application of aloading stress may weaken the OR, as e.g. in 0.15C–1.4Mn–0.25Si–B and 0.33C–1.5Mn–2Si–B steels trans-formed at 700uC under uniaxial compression (Suh et al.,2000; Kang et al., 2003). Primary ferrite grains mayshare a common orientation with neighbouring marten-site resulting from quenching (Cabus et al., 2004b) orexhibit MRs far from KS/KS (Dey et al., 2005).
Intragranular ferrite generally nucleates at inclusions,so that the austenite to ferrite OR strongly depends onthe nature of inclusions. V(C,N) precipitates may breakthe austenite–ferrite OR by introducing incoherentinterfaces with one of the ferrite (or more probablyaustenite) phases, leading to ORs (if any) far from theBain zone, e.g. Cho et al. (2002a and 2002b), Miyamoto
7 Retrieving orientation of parent austenite grain (grey
squares) from those of variants of product phase
formed in it (here, bainitic ferrite in low carbon steel).
After Lambert-Perlade et al. (2004a). Labels I, II and III
denote three Bain zones derived from orientation of
austenite
Table 17 Determination of ORs using EBSD between austenite and martensite in steels
Steel composition Martensite phase Investigation method OR Ref.
9Cr–1Mo–Nb–V a9 (ybcc) MR between a9 variants Near KS Nakashima et al., 20010.2C–12Cr–1Mo–V a9 (ybcc) MR between a9 variants KS Dronhofer et al., 20030.6C–1.5Si–1.5Mn a9 (ybcc) OR between retained c and a9 Near KS Regle et al., 20040.6C–1.5Si–1.5Mn a9 (ybcc) OR between retained c and a9 Close to NW and KS Cabus et al., 2004aInterstitial free and plaincarbon steels (0.2 to 0.6C)
a9 (ybcc) MRs between a9 variants Near KS Morito et al., 2003
AISI 301 0.13C austeniticstainless steel
a9 (ybcc) OR between a9 and c Near KS Lee et al., 2005
19.6Mn–3.1Si–2.9Al e (hcp) OR between e and c Burgers Godet et al., 2005
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International Materials Reviews 2007 VOL 52 NO 2 83
et al. (2003), Hernandez et al. (2003), Furuhara et al.(2003). Intragranular ferrite nucleated at deformationbands of strained austenite is also strongly texturedaccording to the local orientation of austenite crystals(Hurley et al., 1999 and 2000; Hurley and Hodgson,2000). Owing to high local misorientation in austenite,no ‘classical’ OR was evidenced between austenite andferrite (Hurley et al., 1999; Kang and Lee, 2004).
Widmanstatten ferrite also exhibits ORs close to theBain zone with the parent austenite grain, e.g. in ironmeteorites (He et al., 2005), although the OR maycontinuously vary along the interphase interfaces (Heet al., 2006).
Acicular and bainitic ferrite
Mechanisms of austenite to bainite phase transforma-tions are still highly controversial (Hillert et al., 2002).Just as similar ORs prevail for both diffusive (e.g.primary ferrite) and displacive (e.g. martensite) decom-position of austenite, they also prevail for acicular ferrite(Fig. 8) and bainitic ferrite. Plate and lath bainite are amultiple scaled microstructure. Constitutive units arereadily studied by TEM; they cluster into packets (or‘sheaves’) whose size is close to the former austenitegrain size, so that the size and morphology of packets isbest studied using either SEM or light optical micro-scopy, and their crystallography using EBSD. Bainitepackets often contain a high density of LABs (e.g.Tsunekage and Tsubakino, 2002; Lambert-Perlade et al.,2004a). In low carbon steels, the austenite phase isgenerally no longer present at room temperature, so thatORs are generally studied by reconstructing the averagecrystal orientation of individual austenite grains(Lambert-Perlade et al., 2004a) (Fig. 7) or by comparingMRs between packets to those calculated using, e.g. KSand NW ORs (Gourgues et al., 2000). Results obtainedwith EBSD (Table 18) generally well agree with thoseobtained with TEM at the same scale. By consideringthe {001} pole figures, e.g. those from Gourgues et al.(2000) and Verlinden et al. (2001) and the high
frequency of LABs in upper bainite packets, one findsthat no unique OR is generally followed: there arepreferred ORs and continuous (although minor innumber) data points between these preferred ORs. Thesame applies for acicular ferrite (Fig. 8c). Such data canonly be obtained by automated analysis over large areas,which are now made possible by EBSD. Thanks to fieldemission gun (FEG)-SEMs, local ORs along bainitelaths can also be checked. In a 0.6C–1.5Si–1.5Mn steel,the orientation gradient along bainitic laths maintainedthe KS OR with neighbouring film like austenite (Regleet al., 2004).
Owing to its limited angular accuracy, EBSD is notthe best tool to precisely determine an OR, for whichTEM based techniques appear best suited. However,FEG-SEMs now provide high spatial resolution, allow-ing local analysis over wide areas and thus goodstatistics. There is hardly a unique OR at roomtemperature, in particular in metals, because localresidual stresses or plastic rotation develop either duringthe phase transformation or during cooling. The inducedcrystal rotation or distortion may modify a genuine ORobeyed at the very beginning of the phase transforma-tion. Therefore, for most applications, the angularaccuracy of EBSD is generally high enough for ORdetermination.
As with other diffraction techniques, ORs may berepresented with pole figures, including higher indexpoles for which known ORs yield easily recognisablepatterns, which can be compared with experimental datacollected from a single parent grain (Nolze, 2004 and2006). ‘Convoluted’ pole figures may be calculated bythe EBSD software to improve accuracy and sharpnessand facilitate such comparisons (Nolze, 2004). TheRodrigues–Frank space is also increasingly used forORs between fcc and bcc phases (He et al., 2005). Eulerangles may also be useful in certain cases, provided thatan uncertainty interval is suitably given (Nolze et al.,2004).
(a) (b) (c)
8 Electron backscatter diffraction maps of primary and acicular ferrite from two adjacent austenite grains; ferrite grains
in OR with a lower and b upper austenite grain; c corresponding {001} pole figures (no unique OR is found between
ferrite and austenite in either case)
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
84 International Materials Reviews 2007 VOL 52 NO 2
Heterogeneous nucleationInvestigation of homogeneous nucleation requires a highdensity of nuclei and high resolution EBSD to get goodobservation conditions. The EBSD analysis of Fenanodendrites during crystallisation of Fe74Si11B14Ni1amorphous powders during isothermal annealing hasbeen reported (Godec et al., 2006). On the other hand,nucleation in solid state phase transformations is mostoften heterogeneous and relies on crystal structure andorientation of both parent and product phases, so thatheterogeneous nucleation has been extensively studiedwith EBSD. This section focuses on nucleation of bothprecipitates and matrix phases.
Heterogeneous nucleation of precipitates
Austenitic and duplex stainless steels suffer fromboundary precipitation of chromium rich phases afterparticular heat treatments or during high temperatureservice. By analysing a number of GBs with EBSD it wasshown in both duplex (Sato et al., 1999; Sato andKokawa, 1999; Dobransky et al., 2004) and austeniticAISI 304 (Zhou et al., 2001; Hong et al., 2001) stainlesssteels that secondary Cr23C6 and s phases hardlynucleated, or at least nucleated much more slowly atcoincident GBs (or near KS interphase interfaces).Comparing EBSD maps and corrosion tests at the samescale yielded statistically reliable data upon ‘sensitised’grain or phase boundaries. Other criteria, such as planematching boundaries or near coincident axial directiondid not seem to be selective enough. Even the wellknown Brandon criterion for coincident site latticeboundaries was not always selective enough, especiallyfor LABs (Zhou et al., 2001).
Precipitate nucleation has been studied with EBSD ina variety of materials for nuclear power generation. InNi–16Cr–9Fe–xC alloys, random GBs are favouritepaths for intergranular stress corrosion cracking andalso for precipitation of M23C6 (M,Cr) carbides(Alexandreanu et al., 2001). Here again, the GBstructure seems to influence the precipitation kinetics,morphology and size distribution of carbides (Liu et al.,1995). Hydridation of uranium and zirconium alloys isalso affected by the local GB structure; in as casturanium, hydrides form at twin boundaries and atLABs, which are readily imaged with EBSD but notwith light optical microscopy (Bingert et al., 2004).
The spheroidisation of pearlite in 0.36C–0.53Mn–0.22Si steel was further investigated with EBSD, bymonitoring LABs after thermomechanical treatmentsand imaging dissolution and reprecipitation of cementitecoupled with continuous recrystallisation of ferritegrains (Storojeva et al., 2004).
Discontinuous reactions have also been investigatedwith EBSD. In Fe–50 at.-%Co, discontinuous orderingstarts from GBs, but LABs and some coincident sitelattice boundaries seem much less sensitive (Bischoffet al., 1998; Semenov et al., 1998). The local extent ofthe reaction depends on both crystal misorientationand orientation of the GB plane (Bischoff et al., 1998).Such information can be gathered with EBSD (asso-ciated with determination of the local 3D boundarygeometry) over a statistically significant number ofgrains or GBs (more than 200 in this particularexample). This would not be possible by using otherdiffraction methods.T
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oA
uste
mp
ering
Ausfe
rrite
(az
c)C
lose
toK
SM
arr
ow
and
Cetinel,
2000
Cast
iron
3. 6
C–0. 3
Cu–2. 5
Si–
0. 6
Mn–0. 1
5M
oA
uste
mp
ering
Ausfe
rrite
(az
c),
5u
from
KS
Marr
ow
et
al.,
2001
TR
IPste
el
0. 6
C–1. 5
Si–
1. 5
Mn
Ste
pq
uenchin
gB
ain
itez
reta
ined
auste
nite
Locally
rath
er
clo
se
toK
S(1u
from
neig
hb
ouring
pix
els
)R
eg
leet
al.,
2004;
Cab
us
et
al.,
2004a;
Cab
us,
2005
TR
IPste
el
Low
allo
y,
Si–
bearing
ste
el
Hot
rolli
ng
zste
pq
uenchin
g(4
00uC
)B
ain
itez
reta
ined
auste
nite
Not
exactly
KS
(bain
ite
isd
efo
rmed
)G
od
et
et
al.,
2001
TR
IPste
el
0. 4
C–1. 5
Si–
1. 5
Mn
Ste
pq
uenchin
g(4
00uC
)B
ain
itez
15–20%
reta
ined
auste
nite
Not
exactly
KS
God
et
et
al.,
2004
TR
IPste
el
0. 3
9C
–1. 3
7S
i–1. 4
5M
nS
tep
quenchin
g(4
00uC
)B
ain
itez
reta
ined
auste
nite
Clo
se
toK
Sor
NW
Verlin
den
et
al.,
2001
TR
IPste
el
0. 2
C–0. 5
Si–
1. 4
Mn–0. 7
Al
Cold
rolli
ng
zin
terc
riticalte
mp
ering
zste
pq
uenchin
gV
ery
clo
se
toK
S(b
oth
with
TE
Mand
EB
SD
)Z
aeff
ere
ret
al.,
2004
HS
LA
ste
el
0. 0
7C
–0. 3
Si–
1. 5
Mn–C
r–M
o–N
i–N
bTherm
ally
sim
ula
ted
coars
eg
rain
ed
HA
ZU
pp
er
bain
ite
Clo
ser
toN
Wth
an
toK
SG
ourg
ues
et
al.,
2000
Low
carb
on
ste
el
0. 0
7C
–0. 2
Si–
1. 6
Mn–C
r–N
i–N
b–Ti–
V–B
Weld
meta
ld
ep
osit
Acic
ula
rfe
rrite
Clo
ser
toN
Wth
an
toK
SG
ourg
ues
et
al.,
2000
Low
allo
yste
el
0. 2
C–0. 2
Si–
1. 4
Mn–C
r–M
o–N
i(A
533)
Contr
olle
dhot
rolli
ng
Up
per
bain
ite
Clo
ser
toN
Wth
an
toK
SG
ourg
ues
et
al.,
2000
HS
LA
ste
el
0. 0
7C
–0. 3
Si–
1. 5
Mn–C
r–M
o–N
i–N
bTherm
ally
sim
ula
ted
coars
eg
rain
ed
HA
ZU
pp
er
bain
ite
Not
necessarily
KS
or
NW
,als
oclo
se
tooth
er
less
cla
ssic
alO
Rs
Lam
bert
-Perlad
eet
al.,
2004a
Low
carb
on
ste
el
0. 0
9C
–0. 5
Si–
1. 9
Mn–B
–O
Weld
meta
ld
ep
osit
Acic
ula
rfe
rrite
Near
KS
(poin
tanaly
sis
)K
luken
et
al.,
1991
Low
carb
on
ste
els
0. 0
6to
0. 1
C–M
n–S
i–C
u–N
i–C
rS
tep
quenchin
gB
ain
ite
No
MR
betw
een
20
and
47u,
consis
tent
with
KS
or
NW
Dıa
z-F
uente
set
al.,
2003
Low
carb
on
ste
el
0. 2
C–2M
n–0. 3
V–N
–S
Ste
pq
uenchin
gA
cic
ula
rfe
rrite
MR
s,
5u
from
KS
/KS
variants
Miy
am
oto
et
al.,
2003
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 85
Similar studies have been carried out for discontin-uous (cellular) precipitation in Mg–10 wt-%Al (Mg zMg17Al12) (Bradai et al., 2002) and nickel base singlecrystal superalloy (c z c9 z topologically close packedphases) (Lavigne et al., 2004). Here, LABs (Lavigneet al., 2004) or GBs with low index misorientation axes(Bradai et al., 2002) were less sensitive to the involvedreaction. In Cu–4 wt-%Ti, Mangan and Shiflet (1997)combined EBSD with careful serial sectioning to showcomplex interactions between the moving boundary of(fcc azorthorhombic Cu4Ti) colonies and the twinboundaries of the starting microstructure.
Heterogeneous nucleation in matrix phase transformations
Nucleation of ‘major’ phases is of course also stronglyinfluenced by local crystallography. In white cast irons,phase transformation from eutectic cementite (Fe3C)into graphite involved nucleation of graphite at curvedferrite/Fe3C interfaces; however, no nucleation occurredat straight interfaces parallel to the [001] direction ofFe3C and to the normal direction of the product (Songet al., 2002).
Most reported EBSD studies on this topic focused onsteels. In situ light optical microscopy observations offerrite to austenite phase transformation of a coarsegrained Fe–4.8 at.-%Cr–0.008C alloy coated with SiO2
showed preferential nucleation of austenite at triplejunctions involving three random ferrite HABs; thehigher the number of ‘special’ boundaries connected tothe triple junction, the lower the probability foraustenite to be found there (Watanabe et al., 2005b).However, the use of coarse grained material (to get onlyone grain in thickness) does not prevent from freesurface effects and from rather low statistics of results.In a low carbon steel wrapped into an austenitic stainlesssteel foil and heat treated in molten salt baths, nopreferential nucleation of austenite was found near‘special’ grain boundaries (Bruckner et al., 2001).Nucleation of primary ferrite at ferrite/austenite phaseboundaries of a 0.15C–1.4Mn–0.25Si–0.006B steel at700uC was strongly favoured by an applied stress, due toa less strictly observed OR and thus to an increase in thephase boundary energy; this effect seemed stronger thanat austenite GBs (Suh et al., 2000).
Nucleation of intragranular ferrite at secondaryphases or inclusion particles has been widely studiedby TEM (although the probability to find a nucleationsite in a TEM thin foil is low) and by EBSD over a largevariety of inclusions (van der Eijk et al., 1999). Theeffect of deformation of the austenite phase on nuclea-tion of ferrite has also been studied with EBSD,especially in ultrafine grained ferrite formed at deforma-tion bands and cell boundaries of hot rolled austenite(Hurley et al., 2000, Hurley and Hodgson 2001). Incontrast, austenite LABs did not influence markedly itstransformation into martensite after accumulated rollbonding (Kitahara et al., 2004). Nucleation of acicularferrite at TiN particles may occur with an OR betweenevery individual ferrite phase and the TiN particle (Jinet al., 2003). In low carbon steels, the competitionbetween acicular ferrite and bainite at the austenite/ferrite phase boundaries depends on both the chemicalcomposition and the austenite grain size (Dıaz-Fuenteset al., 2003).
The formation of Widmanstatten ferrite sideplates atprimary ferrite crystals already formed at former
austenite GBs was shown thanks to EBSD and serialsectioning to occur not by interfacial instability but bynucleation and growth of individual plates. A distinctGB (and a 5–10u low angle MR) was found with EBSDbetween Widmanstatten sideplates and primary ferritefrom which they were growing (Spanos and Hall, 1996;Phelan and Dippenaar, 2004; Phelan et al., 2005).
As nucleation events are distributed everywherewithin the material, 3D observations are often necessary;TEM seems best suited to fine scale and finely dispersednucleation events; for scarcely distributed nucleationsites, 3D investigations involving EBSD can be bettersuited due to both statistical relevance of data and thepossibility to perform serial sectioning.
Evidence of variant selectionGiven an OR between parent and product phases, thereare a fixed number of crystal orientations of the productphase that can form from a given crystal of the parentphase. These are called ‘variants’. Not all possiblevariants are generally observed in metallographicsections of former individual grains of the parent phase.This can be due to either sampling effects or to a ‘true’variant selection, or both. Variant selection is illustratedin Fig. 9. Variant selection is readily studied at the scaleof the former parent grains by TEM or EBSD and, froma statistical point of view, using the average texture ofthe material (obtained by, e.g. neutron diffraction, XRDor EBSD). Both approaches have successfully usedEBSD in a number of materials. At the scale ofindividual parent grains, the microtexture may readily
(a)
(b)
(c)
a no variant selection; b local, but no average variantselection; c local and average variant selection
9 Two-dimensional view of variant selection (one variant
per shape)
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
86 International Materials Reviews 2007 VOL 52 NO 2
be investigated; at the global scale, such as with XRD,the experimental ODF is compared to that calculatedfrom the orientation of the parent phase by assuming anOR to be strictly obeyed (with or without variantselection criteria).
Martensitic transformations in shape memory and Fe–Nialloys
In Ni–Ti shape memory alloys (SMAs) fabricated byaccumulative roll bonding of Ni and Ti foils, NiTi doesnot directly form from Ni and Ti phases and intermediatecompounds such as Ni3Ti and Ti2Ni are involved.However, a KS OR is observed between Ni and NiTiand a Burgers OR is observed between Ti and NiTi,suggesting that intermediate phase transformationsinvolve variant selection (Inoue et al., 2003). Localvariant selection was shown in Fe–27.5Ni–17.7Co–3.8Tiquenched at –196uC (Bruckner et al., 1999). Variantselection was also suggested by average texture calcula-tions in undeformed (but not in deformed) Fe–28Ni–0.02C (wt-%) (Kestens et al., 2003) (with the Bain OR,however, which is not exactly relevant in that case). Thedetailed microtexture of fork and spear like martensite inCu–7.3Al–8.5Mn (wt-%) (among 24 possible variants of18R1 martensite) showed two groups of two variants ofcommon spatial but different crystal orientations (Wanget al., 2002). Self-accommodating variants, twin relatedby {121}2H and {101}2H planes respectively were found inspear like and fork like martensite of Cu–12.55Al–4.84Ni(wt-%) after transformation of DO3 austenite (Chen et al.,2000). At the scale of parent grains, not all possible variantsare generally found by EBSD in e.g. Fe–29.6 wt-%Ni(Kitahara et al., 2005) and Co thin films (Hesemannet al., 2001), where variants such that {0001}hcp//{111}fcc
at 20u from the surface of the film are favoured.The formation of lenticular martensite in Fe–(31–
32) at.-%Ni bicrystals was extensively investigated usingEBSD, in particular for symmetric ,211. tilt (Fig. 10)and 90u {211} twist boundaries. In single crystals ofsimilar chemical composition, all 12 possible variantswere observed, even when transformation occurredunder applied stress (Ueda et al., 2001b). For bicrystalshaving a tilt boundary and whatever the applied stress
and misorientation between GB and tensile axis, variantselection favoured martensite crystals having a habitplane parallel to the GB plane and enhancing accom-modation of the transformation strain across the GB(Ueda et al., 2001b). This stands generally for other tiltangles under no applied stress (Ueda et al., 2003a) and inprestrained 90u,211. bicrystals (Ueda et al., 2001a).For 90u {211} bicrystals, no strain accommodation waspossible through variant selection across the GB, so thatthe martensite start temperature was lower. However,variant selection was also found (Ueda et al., 2001b)together with local plastic deformation on the other sideof the GB (Ueda et al., 2001a). As residual stresses(resulting from only partial accommodation of trans-formation strains) were higher in the twist than in the tiltbicrystal, a better memory of the starting austenitecrystal orientation was observed after reverse transfor-mation of the twist bicrystal while many LABs appearedin the new austenite phase of the tilt bicrystal (Uedaet al., 2003b and 2004). Owing to the large size of bicry-stals together with the small size of individual marten-site variants, multiscale EBSD study was a key tool toinvestigate phase transformations in such materials.
Non-ferrous metal alloys
Some examples of variant selection evidenced (or notevidenced) by EBSD in titanium and zirconium alloysare given in Table 19 at various scales.
In hot rolled b (bcc) heat treated Cu–40Zn alloys,which then partially transformed into fcc a duringcooling, coarse a grains with strong internal misorienta-tions were found. The a phase formed from cold rolledand annealed (but not recrystallised) b showed only oneBain zone (i.e. near KS or NW ORs clustered aroundthe same 45u,100.b Bain OR). This suggests strongvariant selection that, surprisingly, decreased with thetransformation temperature (Yasuda et al., 1999). Suchvariant selection was also found for (azb) heat treatedCu–42 wt-%Zn, at least for a grains that were clusteredinto bands (Sakata et al., 2000).
In Al–36Mo–17Ti (at.-%), thermal cycling across theb (A2) into Al3Ti (DO22)zMo3Al (A15) eutectoid phasetransformation temperature tended to maintain theinitial orientation of the b phase, showing a memoryeffect on texture and thus involving variant selectionin both eutectoid and reverse phase transformations(Miura et al., 2004 and 2005a). In near-c Ti–46.8Al–1.7Cr–1.8Nb (at.-%), orientation of bothWidmanstatten and lamellar c crystals strongly dependson twinning in the parent a phase, leading to particularMRs between neighbouring Widmanstatten and lamel-lar c phases (Dey et al., 2005).
Steels and ferrous alloys
Variant selection in ferrite (a) to austenite (c) phasetransformation has only been scarcely studied in lowcarbon steels, as the c phase cannot be retained at roomtemperature. In a cold rolled 0.065C–1Mn steel, Parket al. (2002) found variant selection from average texturemeasurements. By measuring average textures of inter-stitial free steels before and after complete a to c thenback to a phase transformations, Ryde et al. (1999)showed no variant selection in the absence of titanium,but a strong memory of initial texture in Ti bearingsteels, possibly due to some retained ferrite in thevicinity of TiN particles. However, such studies are very
10 Variant selection evidenced with EBSD in Fe–
32 at.-%Ni bicrystal containing 90u ,211. tilt bound-
ary and transformed under tension at 45u from tilt
axis. Only variants V29 and V3 are found in each aus-
tenite grain near GB. After Ueda et al. (2001b)
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 87
Tab
le19
Evid
en
ce
of
vari
an
tsele
cti
on
aid
ed
by
EB
SD
inh
cp
meta
lsan
dm
eta
lall
oys
(ite
ms
init
ali
cs
mo
resp
ecifi
call
yad
dre
ss
vari
an
tsele
cti
on
at
pare
nt
gra
inb
ou
nd
ari
es)
Mate
rial
Pare
nt
ph
ase
Pro
du
ct
ph
ase
Scale
of
EB
SD
investi
gati
on
Resu
lts
Ref.
Cold
rolle
dTi
ba
Avera
ge
and
localte
xtu
reIn
ad
ditio
nto
str
eng
thenin
gof
textu
reof
the
bp
hase,
variants
‘alm
ost
com
mon’to
ad
jacent
bg
rain
sare
sele
cte
dG
ey
and
Hum
bert
,2002
Ti
ba9
(lath
mart
ensite)
One
pare
nt
bg
rain
Fro
mM
Rs
betw
een
variants
,th
ere
isvariant
sele
ction
Wang
et
al.,
2003
Ti
ab
(pla
tes
zallo
trio
morp
hs)
Pare
nt
ag
rain
sV
ariant
sele
ction
pro
bab
lym
ain
lyd
uring
gro
wth
of
the
bp
hase
Sew
ard
et
al.,
2004
Cold
rolle
dz
bheat
treate
dIM
I834
Tiallo
yb
Prim
ary
and
second
ary
aM
acro
zones
(larg
eclu
ste
rsof
bg
rain
sof
sim
ilar
cry
sta
lorienta
tion)
No
part
icula
rvariant
sele
ction
for
prim
ary
a;
str
ong
variant
sele
ction
(only
one
textu
recom
ponent)
for
second
ary
aG
erm
ain
et
al.,
2005a
and
2005b
Ti–
6A
l–4V
ba
pla
tes
(diffu
sio
nal
mechanis
m)
Avera
ge
textu
reand
ind
ivid
ual
pare
nt
gra
ins
No
variant
sele
ction
inavera
ge;
at
least
thre
efa
mili
es
of
ap
late
sp
er
bg
rain
Hum
bert
et
al.,
1994
Up
to30%
hot
rolle
dTi–
6A
l–4V
ba
(diffu
sio
nal)
Avera
ge
textu
reand
ind
ivid
ual
pare
nt
gra
ins
No
variant
sele
ction
inavera
ge;
all
12
variants
found
ineach
pare
nt
bg
rain
Gey
et
al.,
1996
Hot
rolle
dz
heat–
treate
dTi–
6A
l–4V
ba
(diffu
sio
nal)
Avera
ge
textu
reand
alo
ng
ind
ivid
ualG
Bs
Evid
ence
of
variant
sele
ction
at
GB
stw
oaltern
ating
variants
of
prim
ary
a;
LA
Bb
etw
een
Wid
mansta
tten
and
prim
ary
aif
no
sp
ecia
lO
Rb
etw
een
bg
rain
s;
,0001
.a//
com
mon
,110
.b
ifexis
ting
Sta
nfo
rdand
Bate
,2004
Ti–
6A
l–2S
n–4Z
r–6M
ob
Wid
mansta
tten
and
GBa
Pare
nt
bg
rain
sW
idm
ansta
tten
variants
as
clo
se
as
possib
leto
prim
ary
aw
hile
ob
eyin
gth
eB
urg
ers
OR
with
the
pare
ntb
gra
ins
into
whic
hth
ey
are
gro
win
g.
Ifnot
possib
le,
para
llelclo
se
packed
directions
as
clo
se
as
possib
leto
the
pare
nt
GB
Bhattachary
ya
et
al.,
2003
Ti–
8A
l–xV
(laser
dep
osited
)b
GB
aA
long
one
hug
eG
Bover
the
whole
com
positio
nra
ng
eTw
oaltern
ating
variants
with
{0001} a
//to
the
sam
e{1
10} b
,m
isoriente
d10u
from
anoth
er;
OR
str
ictly
ob
eyed
on
one
sid
e,
as
clo
sely
as
possib
lew
ith
the
oth
er
pare
nt
gra
in
Banerjee
et
al.,
2004
Zircalo
y–4
ba
(diffu
sio
nal)
Avera
ge
textu
reThere
isvariant
sele
ction,
esp
ecia
llyw
hen
no
str
ess
isap
plie
dG
ey
et
al.,
2002
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
88 International Materials Reviews 2007 VOL 52 NO 2
difficult and the majority of results concern transforma-tions from austenite to ferrite (or martensite) in steels.
Diffusional formation of primary ferrite
Globular ferrite growing from residual d ferrite of stripcast 0.1C–0.17Si–0.6Mn–P steels obviously shows astrong texture (with only internal LABs) within align-ments (Umezawa et al., 2003). If no d ferrite is retained(i.e. for the majority of ferritic steels), variant selectiondepends on the deformation state of parent austenite. Inweakly prestrained austenite of a 0.3Si–1.5Mn steel,coarse primary ferrite formed from coalescence of agrains of close crystal orientations (i.e. from the samevariant). For strains higher than y40%, neighbouring agrains formed as misoriented variants (Torizuka et al.,2000). In a 0.002C–1.66Mn–Si–Mo–Ni–Al steel vacuumcast into 2 mm thick sheets and then hot rolled, a strongtexture was found in individual former c grains, stronglydepending on the c crystal orientation, but wellpredicted by using all 24 KS variants, so that no variantselection could be evidenced (Hurley et al., 1999).
The diffusional formation of primary ferrite may alsobe significantly driven by free surface energy, such as incoarse grained ferritic steels obtained by intercriticalannealing (i.e. in the azc temperature range) in adecarburising atmosphere. Carbon depletion startingfrom free surfaces induced transformation of the twophase alloy into a fully ferritic microstructure.Abnormal grain growth then occurred due to the lowerfree surface energy of {100} grains and to some memoryof the initial texture and grain size effect, so that a coarsegrained microstructure with a desirable {100} texturewas finally obtained (Tomida, 2003; Tomida et al., 2003;Dzubinsky et al., 2004).
Transformation of austenite into Widmanstatten ferrite,bainite and martensite
Whatever the transformation mechanism, variant selec-tion in bainite and Widmanstatten ferrite has alreadybeen widely investigated with EBSD (Table 20). Owingto the coarse size of bainite packets with respect to theparent c grain size, a relatively low number of packetsare generally observed within a given c grain by 2Dsectioning. This leads to significant sampling effects. Thecomplex, non-convex and non-equiaxed shape of bainitepackets may also lead, even in the absence of variantselection, to an inhomogeneity in variant distributiondue to the morphological orientation of packets withrespect to the sample surface plane (Fig. 11). This purelygeometric phenomenon has recently been called ‘pseudovariant selection’ (Cabus, 2005). By assuming a givenshape and habit plane of bainite packets, one canattempt to process microtexture data to take pseudovariant selection into account in each parent austenitegrain (Cabus, 2005). Nevertheless, the vast majority ofliterature data does not take either these considerationsor even the (sometimes low) number of analysed c grainsinto account.
Variant selection is generally investigated at the scaleof the former c grain (i.e. from the resulting local textureor number of bainite variants). In some instances,however, such as in fatigue and cleavage crackpropagation, the relevant parameter is the misorienta-tion between neighbouring packets only. Whatever thebainitic or lath martensite microstructure, no misor-ientation angle of y20u was found between neighbour
packets, although it could possibly exist according tonear KS ORs (Fig. 12). Thus, as evidenced by localhistograms of misorientation angles, there are ‘forbid-den’ pairs of neighbours and thus ‘local’ variantselection, even if it may have no consequence on theaverage texture (Gourgues et al., 2000; Gourgues, 2003).
In hcp e martensite (Godet et al., 2005) showed onlytwo over the four possible {111}c parallel to {0001}e in a19.6Mn–3.1Si–2.9Al steel, suggesting that there couldalso be variant selection in this case. In AISI 304austenitic stainless steel tensile deformed at –60uC, emartensite appeared in parent grains as families ofparallel bands (one variant per family). Within thesebands, e martensite then probably partially transformedinto bcc a9 martensite having one {110} plane parallel tothe (0001)e plane. Six variants out of 24 were thusselected for each band family. Moreover, there wasfurther local variant selection among these six variants(Gey et al., 2005).
Variant selection at parent grain boundaries
Intergranular nucleation of a new matrix phase involvesat least three contributions to the change in free energy:
(i) GB interfacial energy
(ii) chemical variation in free energy due to phasetransformation
(iii) interface and strain energy of coupled new phaseand individual matrix grains.
(i) and (iii) are variant dependent, so that new variantsmay tend to obey a given OR with both parent grains.This is best studied with EBSD, especially if the parentgrain size is coarse. Examples from hexagonal metalsand metal alloys are given in italics in Table 19.
In Cu–40Zn with {001},110. textured bcc b, fcc avariants either of close crystal orientations or in twinMR formed on either side of the b GBs. Coarse acolonies nucleated on either side of the same GBgenerally shared at least a common low index direction,with ,111.a directions close to the common ,110.b.In contrast to Ti alloys, this has only little effect on theaverage texture (Stanford and Bate, 2005). In Ni–43Crheat treated in the fcc c phase range and then cooleddown slightly below the solvus of the bcc a phase, whenexisting, the only variant selected was that having a KSOR with both c grains; otherwise, selected variantsobeyed the KS OR with one c grain, and had their closepacked plane of the OR as close as possible to the c GBplane (Adachi and Tsuzaki, 2005; Adachi et al., 2005).
In cold rolled low alloy (0.065C–0.05Si–0.99Mn) steelannealed in the c temperature range, EBSD resultssuggested that a memory of the initial a crystalorientation still existed, with selection of a variants inKS OR with the neighbouring c grains (frequently in KSOR themselves with the same initial a crystal) (Brucknerand Gottstein, 2001). Such variant selection was alsoobserved in a 0.15C–0.3Si–1.42Mn steel after hot torsion(Novillo et al., 2004). In austenitic 0.0015C–22Cr–(8–9)Ni stainless steels, gas tungsten arc welding followedby liquid tin quenching showed that at least at thebeginning of bcc d to fcc c phase transformation, cvariants at d GBs were more frequently in KS ORwith both neighbouring d grains than with one d grainonly (Inoue et al., 1998). The same result was foundin annealed and quenched duplex stainless steels(Gourgues et al., 2004; Monlevade et al., 2006) (Fig. 13).
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 89
Tab
le20
Evid
en
ce
of
vari
an
tsele
cti
on
aid
ed
by
EB
SD
info
rmati
on
of
bain
itic
or
Wid
man
sta
tten
ferr
ite*
Allo
yco
mp
osit
ion
Fin
al
mic
rostr
uctu
reA
naly
sis
tech
niq
ue
Resu
lts
Vari
an
tsele
cti
on
Ref.
3. 2
C–1C
u–0. 5
5M
n–2S
iA
scast
gre
yiron:
bain
itez
reta
ined
cN
um
ber
of
variants
per
cg
rain
Severa
lvariants
per
cg
rain
Not
enoug
hd
ata
toconclu
de
Ferr
yand
Xu,
2004
Ni–
Fe-r
ich
mete
oritic
min
era
lW
idm
ansta
tten
aO
ne
hug
ec
gra
inanaly
sed
All
12
near
NW
variants
were
found
No
He
et
al.,
2006
Low
allo
y,
Sib
earing
(TR
IP)
Bain
itez
reta
ined
cIn
div
idualorienta
tion
of
variants
Variants
sele
cte
daccord
ing
toactive
slip
syste
ms
Yes
God
et
et
al.,
2001
0. 3
9C
–1. 3
7S
i–1. 4
5M
n(T
RIP
)Lath
bain
itez
reta
ined
cN
um
ber
of
morp
holo
gic
alvariants
;avera
ge
textu
rew
ithin
pare
nt
cg
rain
Low
er
num
ber
of
morp
holo
gic
alvariants
for
ste
pq
uenchin
gfr
om
defo
rmed
c;m
axim
um
inte
nsity
of
textu
revaries
from
one
cg
rain
toanoth
er
Verlin
den
et
al.,
2001
0. 1
9C
–1. 4
6S
i–1. 5
7M
nand
0. 3
1C
–0. 3
4S
i–1. 5
7M
n–1. 2
3A
l(T
RIP
)B
ain
itez
reta
ined
cA
vera
ge
textu
reof
ccalc
ula
ted
from
that
of
bain
ite
{001},
100
.c
weaker
(resp
ectively
str
ong
er)
than
pre
dic
ted
for
Al–
Si(r
esp
ectively
Alfr
ee)
ste
el
Possib
lyyes
De
Meyer
et
al.,
2001
0. 6
C–1. 5
Si–
1. 5
Mn
(TR
IP)
Bain
ite
z.
20%
czm
art
ensite
Num
ber
of
variants
per
cg
rain
(over
afe
wc
gra
ins)
Less
than
the
24
possib
leK
Svariants
;clo
se
packed
pla
nes
of
KS
OR
are
para
llelto
only
one
or
two
{111} c
pla
nes;
possib
lesam
plin
geff
ects
?
Possib
lyyes
Reg
leet
al.,
2004;
Cab
us
et
al.,
2004a
0. 7
9C
–1. 5
6S
i–1. 9
8M
n–1. 0
1A
l–0. 2
4M
o–1. 0
1C
r–1. 5
1C
o(T
RIP
)C
arb
ide
free
bain
ite
Tra
ce
of
bain
ite
packets
,m
orp
holo
gy
of
cry
sta
llog
rap
hic
packets
Tra
ces
are
para
llelto
maxim
alshear
pla
nes;
packets
are
bett
er
alig
ned
iftr
ansfo
rmed
und
er
com
pre
ssio
n
Yes
Hase
et
al.,
2004
0. 6
C–1. 5
Si–
1. 5
Mn
(TR
IP)
Bain
itez
.20%
czm
art
ensite
Avera
ge
aand
cte
xtu
res
takin
gp
seud
ovariant
sele
ction
into
account
Clo
se
packed
pla
nes
of
near
KS
OR
are
para
llelto
one
or
two
dom
inating
{111} c
pla
nes.
Yes
Cab
us,
2005
0. 4
C–1. 5
Si–
1. 5
Mn
(TR
IP)
Bain
itez
15–20%
cfr
om
cd
efo
rmed
by
0. 2
or
0. 8
Num
ber
of
KS
variants
per
pare
nt
cg
rain
All
24
variants
are
found
for
e50. 2
,not
all
for
e50. 8
;in
tern
alm
isorienta
tion
incre
ases
with
incre
asin
gd
efo
rmation
of
auste
nite
Yes
God
et
et
al.,
2004
0. 2
2C
–1. 6
Si–
1. 5
Mn–0. 0
45N
b(T
RIP
)B
ain
itez
reta
ined
cN
Wvariants
v.
active
part
iald
islo
cation
slip
syste
ms
Ina
giv
en
cg
rain
,b
oth
variants
associa
ted
toactive
slip
syste
mand
op
posite
syste
mare
found
Possib
lyyes
Jonas
et
al.,
2005
0. 0
4C
–0. 2
Si–
1. 5
Mn–N
i–N
b(H
SLA
)E
long
ate
db
ain
ite
alo
ng
the
hot
rolli
ng
direction
MR
sb
etw
een
bain
ite
‘colo
nie
s’
Low
(,15u)
mis
orienta
tions:
one
colo
ny
per
cg
rain
.P
ossib
lyone
colo
ny
51
to2
clo
sely
oriente
dK
Svariants
Possib
lyyes
Mats
uoka
et
al.,
1999
0. 0
7C
–0. 3
2S
i–1. 5
Mn–C
r–M
o–N
b–V
(HS
LA
)Lath
bain
ite
from
coars
eg
rain
ed
cM
Rs
betw
een
packets
com
pare
dto
KS
or
NW
rela
ted
MR
sC
lose
packed
pla
nes
of
OR
para
llelto
diffe
rent
{111} c
pla
nes
inneig
hb
ouring
packets
;th
ere
are
locally
forb
idd
en
MR
s
Locally
yes
Gourg
ues
et
al.,
2000;
Lam
bert
-Perlad
eet
al.,
2004a
0. 2
C–0. 2
5S
i–1. 3
8M
n–C
r–M
o–N
i–A
l–N
b–Ti–
V(A
533)
Tem
pere
db
ain
ite
MR
sb
etw
een
packets
com
pare
dto
KS
or
NW
rela
ted
MR
sU
pp
er
bain
ite:
clo
se
packed
pla
nes
of
OR
para
llel
tod
iffe
rent
{111} c
pla
nes
inneig
hb
ouring
packets
;lo
wer
bain
ite:
clo
se
tom
art
ensite
mic
rote
xtu
re(s
ee
section
on
mic
rote
xtu
res)
Locally
yes
Gourg
ues
et
al.,
2000
0. 0
5C
–0. 5
Si–
1. 6
Mn–V
(HS
LA
)B
ain
itez
reta
ined
cA
vera
ge
textu
re,
NW
OR
sE
xp
erim
enta
lte
xtu
renot
the
sam
eas
calc
ula
ted
one
Yes
Hum
bert
et
al.,
2002b
0. 1
5C
–2. 2
5C
r–1M
o,
0. 0
5C
–9N
i,0. 0
7C
–0. 4
Si–
1. 4
Mn–V
Up
per
bain
ite
(sim
ula
ted
coars
eg
rain
ed
HA
Zs)
MR
sb
etw
een
packets
,siz
eof
cle
avag
efa
cets
Up
per
bain
ite:
many
inte
rnalLA
Bs;
low
er
bain
ite:
many
HA
Bs;
there
are
locally
forb
idd
en
MR
s.
Cle
avag
efa
cets
of
sam
esiz
eas
cg
rain
s
Locally
yes
Gourg
ues,
2003
0. 2
C–0. 5
Si–
1. 4
Mn–0. 7
Al(T
RIP
)Ferr
itez
bain
ite
shellz
reta
ined
cLocalM
Rs
LA
Bs
betw
een
ferr
ite
and
bain
ite:
bain
ite
orienta
tion
dic
tate
db
yth
at
of
ferr
ite
Yes
Zaeff
ere
ret
al.,
2004
*TR
IP:
transfo
rmation
ind
uced
pla
sticity
aid
ed
ste
els
;H
SLA
:hig
hstr
eng
thlo
wallo
y.
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
90 International Materials Reviews 2007 VOL 52 NO 2
In summary, EBSD investigations of various metalalloy systems showed that at GBs of the parent phase,variant selection may occur, favouring those variantseither obeying as strictly as possible the OR with bothparent grains, or having a special OR with the GB plane.
Examples of practical applications of variantselectionGrain boundary engineering (GBE) through phasetransformations
Grain boundary engineering is a method by which,through well controlled metallurgical and thermome-chanical processing, ‘special’ GBs are strongly favouredleading to improved product properties (see e.g.Watanabe, 1984). This has been made possible byextensive EBSD characterisation (Gourgues, 2002).The GBE is generally achieved through series of hotworking and annealing treatments, making use of grainboundary motion, recrystallisation and annealing twin-ning. However, variant selection in phase transforma-tions could be another way to produce ‘special’ GBs
between grains of product phases through controlledphase transformations. The feasibility of GBE wasshown in orthorhombic Ti–Al–Nb alloys by variousmetallurgical routes (Li et al., 2004; Boehlert et al., 2004;Li and Boehlert, 2005a). According to the DO19 a2 orordered B2 parent phases, special MRs were foundbetween orthorhombic O grains. Some MRs betweenvariants were favoured (e.g. 65u,001. or 90u misor-ientations). Other MRs seemed to be locally forbidden(e.g. 30u,001. starting from the a2 phase). Localtexture could thus possibly be tailored through variantselection, even for random average textures.
The GBE might also be achieved by phase transfor-mation under particular conditions, such as in para-magnetic fcc austenite into ferromagnetic (pearlitic) bccferrite in a medium carbon steel (Zhang et al., 2005) andin a 1.0C–Si–Mn–Cr steel (Zhang et al., 2006). A highapplied magnetic induction (12 to 14 T) both increases
11 Sectioning effects inducing pseudo variant selection,
here in case of rod shaped phases: area intercepted by
sectioning plane (in white) is schematically repre-
sented in dark grey; variants having their long axis clo-
sest to sectioning plane have high apparent fraction
12 Histogram of misorientation angles in low carbon
steels, showing apparently ‘forbidden’ pairs of neigh-
bouring ferrite variants. Arrows denote theoretically
possible angles between pairs of variants in KS or
NW ORs with parent austenite grain into which they
grow. After Gourgues et al. (2000)
a b
a light optical micrograph; b EBSD map of same area (ferrite in black, austenite in various grey levels)13 Growth of austenite colonies across parent ferrite grain boundaries in cast duplex stainless steel. Austenite colonies
keep near KS OR with both ferrite grains. Ferrite GB is delineated with broken white line. After Gourgues et al.
(2004)
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 91
the transformation temperature and strongly decreasesthe amount of LABs between ferrite crystals. Grainboundary engineering on one phase (or more) at a timemight also be achieved in two phase alloys such asin azb9 Cu–40 wt-%Zn alloys (Lee et al., 2003).Conversely, variant selection could possibly be used toavoid certain ‘low strength’ GBs in product phase, suchas 90u MRs between hcp a grains in Ti–6Al–4V (Bieleret al., 2005a); this would be very useful to avoidextensive cavitation of such GBs during upset forging.
Preventing growth and coarsening of ferritic steels
Among the practical issues of phase transformations,structure refinement is achieved by monitoring not onlythe number of nucleated grains, but also the evolution ofmicrotexture during or after phase transformation. Asan example, it is still very difficult to achieve a grain sizelower than y5 mm in ferritic steels, although the highnumber of nucleated grains should lead to a finer ferritegrain size. This is attributed to coarsening, a phenom-enon which is intimately related to the microtexture ofthe initially formed ferrite. Thus, coarsening could belimited by better control of the phase transformationstage, where EBSD is of great help as illustrated inTable 21. The EBSD was used here to investigate, over asignificant number of grains, the deviation from strictOR (and thus decreasing probability that ferrite grainsof the same crystal orientation nucleate close to eachother and eventually coalesce) and to quantitativelydefine the grain size by a threshold criterion involvingthe misorientation angle.
Microtexture: relationships betweenmorphology and local crystallographyOwing to anisotropic interfacial energy or to transfor-mation strains, there can be strong correlation betweencrystal orientation and morphology resulting from solidstate phase transformations (see e.g. habit planes ofmartensite or of certain precipitates). The morphologyof phases may be observed using conventional lightmicroscopy or SEM imaging techniques. However, ifproduct phases come in intricate or complex shapedentities, or if only 2D trace analysis is possible, suchtechniques may give a confusing appearance of productphases. One must then distinguish between what appearsby imaging (here denoted as ‘morphological’ entities)and ‘true crystals’ surrounded by user defined bound-aries (here denoted as ‘crystallographic’ entities). Theright scale is generally that of the SEM/EBSD technique.Several cases have been encountered using EBSD,namely:
(i) several morphological units may belong to thesame crystallographic unit, as in e.g. pearlitecolonies of pearlitic steels (Aernoudt et al.,2005), in misoriented intragranular ferritenucleated at a given inclusion, in acicular ferritein certain low carbon steels (Dıaz-Fuentes et al.,2003), in hcp a phase colonies from bcc b phasesin Ti–6Al–2Sn–4Zr–6Mo (azb) titanium alloy[close Burgers variants in that case (seeBhattacharyya et al., 2003)] and in plates orconglomerates of laths in proeutectoid Fe3C ofa Fe–1.34C–13.1Mn steel (Mangan et al., 1999)
(ii) several crystallographic units may be morpho-logically parallel and become undistinguishablefrom each other in 2D metallographic sections.
This is the case of groups of variants inmartensitic Cu–15.4Al–8.9Mn (at.-%) SMA(Wang et al., 2002). This is also the case for‘blocks’ of lath martensite or even bainite insteels (see below)
(iii) crystallographic units may correspond to a few,very intricate morphological entities, whichappear as ‘knitted’ or ‘woven’ (Fig. 14a). Thisis clearly the case in upper bainite steelmicrostructures (Gourgues et al., 2000;Nohava et al., 2003; Lambert-Perlade et al.,2004a) and also possibly (by closely looking atmicrographs) in hcp a colonies in Ti–5Ta–1.8Nb (Karthikeyan et al., 2005). This micro-texture appears when variants of close crystalorientations, but highly misoriented in mor-phology are selected. Although there are in factmany crystals in such a crystallographic entity,these are separated by LABs only, which do notstrongly affect properties such as resistance tocleavage or fatigue crack propagation. Thisleads to coarser unit crack path and lowertoughness properties (Gourgues et al., 2000;Kim et al., 2000; Dıaz-Fuentes et al., 2003;Nohava et al., 2003; Lambert-Perlade et al.,2004a). As a result, no obvious correlationbetween morphology and crystal orientationcan be found in this case (Gourgues et al., 2000;Nohava et al., 2003)
(iv) intricate, both crystallographically and mor-phologically misoriented units may be found.This is the case of, e.g. Widmanstatten cemen-tite in Fe–0.8C–12.3Mn steel (Mangan et al.,1997) and in acicular ferrite of steel welddeposits (Gourgues et al., 2000) (Fig. 8)
(v) irregular entities, where no particular correla-tion between morphology and crystal orienta-tion can be evidenced, such as primary hcp aformed from bcc b in Ti–8Al–6V (Banerjee et al.,2004) and individual pearlite colonies in fullypearlitic steels (Aernoudt et al., 2005).
In some instances, coupled information on crystal andmorphological orientations may give information aboutthe phase transformation scheme during metal proces-sing, e.g. plate hcp a plates in extruded Zr–2.5Nbpressure tubes; in this example, some pro-monotectoid aphase already formed before extrusion, whereas discreteclusters of a particles having their ,0001. axes parallelto the tube axis appeared during extrusion only(Griffiths et al., 1998, Holt and Zhao, 2006).
The microtexture of lath martensite in steels mayreadily be determined with EBSD, at least from 2Dmaps, even for coarse parent austenite grains (Fig. 14b)(Nakashima et al., 2001; Morito et al., 2003 and 2006;Dronhofer et al., 2003). The following description ofmartensite microtexture, which was built from EBSDresults only, is based on the KS OR, even though thisOR is not strictly followed in lath martensite. It hasrecently been extended, with some modifications, tobainite in Fe–9Ni steels transformed between 450 and350uC (Furuhara et al., 2006). Laths sharing almost thesame crystal and morphological orientation are gatheredtogether into so called ‘subblocks’. Subblocks having thesame morphological orientation but slightly differentcrystal orientations (same parallel close packed planes
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
92 International Materials Reviews 2007 VOL 52 NO 2
Tab
le21
Ele
ctr
on
backscatt
er
dif
fracti
on
investi
gati
on
of
gro
wth
an
dco
ars
en
ing
of
ferr
itic
ste
els
rela
ted
toau
ste
nit
eto
ferr
ite
ph
ase
tran
sfo
rmati
on
Ste
el
co
mp
osit
ion
Th
erm
om
ech
an
ical
treatm
en
tS
tart
ing
cm
icro
str
uctu
reO
RR
esu
lts
on
co
ars
en
ing
Ref.
0. 1
5C
–0. 2
5S
i–1. 4
Mn–0. 0
06B
1150uC
for
3m
inth
en
10%
unia
xia
lcom
pre
ssio
nstr
ain
at
700uC
,¡
50
sd
well
befo
recoolin
g
Tra
nsfo
rmation
str
ess
free
and
und
er
str
ess
KS
,w
eaker
for
und
er
str
ess
than
for
str
ess
free
Aff
ects
main
lyfe
rrite
nucle
ation
Suh
et
al.,
2000
0. 1
7C
–0. 3
Si–
1. 5
Mn
1200uC
,1
min
then,
com
pre
ssed
at
750uC
zcoolin
gP
lastically
defo
rmed
(20–50%
)N
ot
sp
ecifie
dFerr
ite
gra
ins
more
eq
uia
xed
,H
AB
more
num
ero
us
with
incre
asin
gp
lastic
str
ain
of
c(b
ut
less
than
30
gra
ins
stu
die
din
cert
ain
sam
ple
s)
Torizuka
et
al.,
2000
0. 0
022C
–0. 2
2S
i–1. 6
6M
n–M
o–N
i47%
pla
stic
str
ain
at
850uC
zair
coolin
gH
eavily
defo
rmed
KS
Most
ferr
ite
gra
ins
are
surr
ound
ed
by
HA
Bs
(but
KS
/KS
MR
snot
cle
arly
vis
ible
on
pole
fig
ure
s)
Hurley
et
al.,
1999
0. 1
Cp
lain
carb
on
ste
el
Hot
rolle
dto
str
ain
of
1. 5
¡d
well
for
10
sth
en
wate
rq
uenchin
gN
ot
sp
ecifie
dE
xte
nsiv
ecoale
scence
of
ferr
ite
gra
ins
(LA
Bs
dis
ap
pear)
Kelly
et
al.,
2002
0. 1
7C
–0. 4
5S
i–1. 5
Mn–A
l–V
Hot
tors
ion
Auste
nite
gra
insiz
eof
14
or
84mm
Not
sp
ecifie
dC
oale
scence
of
ferr
ite
during
coolin
g,
contr
olle
db
y2D
(raft
s)
or
3D
meeting
of
ag
rain
s;
EB
SD
giv
es
the
ag
rain
siz
ed
istr
ibution
Bela
diet
al.,
2004
0. 1
5C
–0. 3
Si–
1. 4
2M
n–N
b–V
Hot
tors
ion
(sim
ula
ting
roug
hin
gand
finis
hin
gro
lling
)D
efo
rmed
or
recry
sta
llised
KS
Defo
rmed
c:fe
wLA
Bs
(,15u)
due
tovariant
sele
ction
or
pla
stic
rota
tion;
recry
sta
llised
c:m
any
LA
Bs,
ag
rain
siz
econtr
olle
db
ycoale
scence
favoure
db
ysuch
variant
sele
ction
Novill
oet
al.,
2004
0. 1
5C
–0. 2
5S
i–1. 4
Mn–B
and
0. 3
3C
–2. 1
Si–
1. 5
Mn–B
1200uC
dow
nto
700uC
,10%
com
pre
ssio
nz
coolin
gd
ow
nto
400uC
with
or
without
str
ess
KS
Shift
off
KS
OR
incre
ases
ifth
elo
ad
issusta
ined
,le
ss
coale
scence
of
neig
hb
ouring
ag
rain
s(b
ut
many
OR
s.
15u
from
KS
inth
ep
ap
er)
Kang
et
al.,
2003
0. 0
82C
–0. 3
6S
i–1. 5
Mn–N
b–V
Multip
ass
hot
tors
ion
Heavily
defo
rmed
Not
sp
ecifie
dFerr
ite
gra
ins
meet
firs
tat
cG
Bs;
there
iscoale
scence
behin
dth
etr
ansfo
rmation
front;
EB
SD
giv
es
ferr
ite
fraction
and
siz
ed
istr
ibution
and
mig
rating
inte
rface
are
ap
er
unit
volu
me
Cotr
ina
et
al.,
2004
0. 0
32C
–0. 1
5S
i–0. 7
4M
n–N
b1150uC
for
5m
inth
en
0. 8
com
pre
ssiv
estr
ain
at
845uC
at
various
str
ain
rate
sth
en
wate
rq
uenchin
g
Defo
rmed
Str
ain
ed
at
0. 0
01
s–1:
dynam
icsoft
enin
gvia
str
ain
ind
uced
transfo
rmation,
rand
om
lyoriente
da
gra
ins,
little
coale
scence;
0. 1
s–1:
most
aG
Bs
are
HA
Bs,
finer
final
ag
rain
siz
e
Eg
hb
ali
and
Ab
dalla
h-Z
ad
eh,
2006
0. 1
9C
–1. 5
Si–
1. 5
Mn–A
l–(0
. 003
or
0. 0
1)N
(TR
IP)
Sim
ula
ted
continuous
annealin
gR
ecry
sta
llised
with
more
or
less
num
ero
us
AlN
part
icle
sN
ot
sp
ecifie
dFin
er
ferr
ite
gra
insiz
eand
hig
her
volu
me
fraction
of
reta
ined
c(b
oth
measure
dw
ith
EB
SD
)w
ith
incre
asin
gth
ed
ensity
of
AlN
(measure
dw
ith
TE
M)
Baik
et
al.,
2006
0. 1
5C
–1. 5
Si–
1. 5
Mn
(TR
IP)
950
or
1200uC
then
hot
tors
ion
at
various
tem
pera
ture
sD
efo
rmed
Not
sp
ecifie
dThe
ferr
ite
gra
insiz
e(E
BS
Dm
easure
ments
)d
ep
end
son
the
therm
alm
echanic
altr
eatm
ent;
best
results
ifd
efo
rmed
at
Ar 3
tem
pera
ture
God
et
et
al.,
2006
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 93
but different ,111. directions parallel to the ,110.
close packed directions of parent austenite) are in turnclustered into ‘blocks’, so that particular MRs betweenvariants (e.g. 60u,111. twin MR) are found betweenblocks. These blocks are parts of ‘packets’. Within onepacket, all laths are morphologically parallel; however, apacket consists in a variety (up to six) of intricate, highlymisoriented crystals. Martensite blocks are highlyintricate over the whole martensite packets (Gourgueset al., 2000; Nakashima et al., 2001; Dronhofer et al.,2003) although they may look equiaxed in 2D sections
(Nakashima et al., 2001). No particular neighbourselection was found among the six possible blocks in agiven packet (Morito et al., 2003). In martensite formedfrom fine grained austenite, one packet may dominate(at least in 2D sections) (Morito et al., 2005), or less thansix variants may be found in every packet (Kitaharaet al., 2006). Such results confirm those previouslyobtained with TEM, yet with much higher statisticalsignificance.
The determination of crystallographic units needscareful selection of the threshold angle used in EBSD todefine their boundaries. In fatigue and cleavage sen-sitive materials, this is related to crack arrest, e.g. inaustempered ductile cast iron, whose packet boundariesdefined with EBSD may in fact stop fatigue crackpropagation (Marrow et al., 2001). The size of theseunits can be used to compare steel microstructures suchas bainite and martensite (Tsunekage and Tsubakino,2002). Note, however, that cracks only arrest atcontinuous boundaries, while 2D sections also take intoaccount small packets locally embedded in larger ones.Consequently, the true unit crack path may be muchlarger than that measured from 2D EBSD maps [e.g. inlath martensite steel microstructures (Gourgues, 2003)].Local orientation gradients also disturb the definition ofrelevant boundaries. They are common along bainitelaths [e.g. in high carbon steels (Regle et al., 2004) andalong the LAB network within crystallographic units ofpearlitic steels (Aernoudt et al., 2005)].
Resulting average texturesThanks to the development of high speed systems,EBSD is now increasingly used for average texturemeasurements, avoiding the problem of calculating theODF from pole figures (as for XRD or neutrondiffraction). The EBSD is also used to infer the textureof the parent phase from that of low temperatureproduct phase. Here, spatial information provided byEBSD is very useful. Such ‘texture history’ is ofoutstanding value to understand and improve proces-sing routes to get optimal texture of final product at lowcosts.
Retrieving texture of parent phase
To infer the ODF of the parent phase from that of theproduct phase, one can work directly with individualgrains (i.e. by calculating individual orientations ofparent phase grains from resulting variants of theproduct phase). This is here referred to as a ‘local’approach. Another possibility is to work with averageODFs only. Here, EBSD is of particular use if the parentphase is coarse grained (e.g. b grains hundreds ofmicrons in size in heat treated titanium alloys). This willbe referred to as a ‘global’ approach.
The ‘local’ approach was first used in materials wherethe parent grains were still easily reconstructed withimaging techniques. One can then select appropriatevariants of the product phase, determine their crystalorientation by EBSD point or map analysis and thencalculate the orientation of the parent grain by assumingan OR between parent and product phases. This workswell in titanium and zirconium alloys, where the BurgersOR is rather strictly obeyed. Two or three variants aregenerally needed to calculate every grain orientation ofthe parent b phase (Humbert et al., 1995). This methodwas successfully used for Ti–6Al–4V, a number of
a
b
a partially transformed upper bainite with highlymisoriented sets of parallel units sharing close crystalorientation as shown in {001} pole figure: afterLambert-Perlade et al. (2004); b fully transformed lathmartensite, where each block ‘B’ delineated by whitelines always contains same pair of low misoriented var-iants ‘V’: after Morito et al. (2003)
14 Microtexture evidenced using EBSD in product phase
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
94 International Materials Reviews 2007 VOL 52 NO 2
triplets of a plates giving about 100 orientations offormer b grains (Humbert et al., 1994; Stanford andBate, 2004) and by using both primary and colony aphase in equiaxed b grains (Humbert et al., 1996;Moustahfid et al., 1997a) and coarse a colonies of pureTi (Gey and Humbert, 2002). The method requires,however, that GBs of the parent phase be readilyrecognised. It works well even if there is variantselection, but also requires that the OR be strictlyobeyed. This is not the case in steels, where no uniqueOR is generally found; the uncertainty in the ORstrongly decreases the accuracy of results concerning theparent phase: a departure of .5u from the OR preventsfrom finding the right orientation of the parent grain(Decocker et al., 2003). In steels, provided that there areenough ferrite crystals formed within a former austenitegrain, it is possible to determine the orientation of thatformer c grain by an iterative graphical method using{001} pole figures (Gourgues et al., 2000; Cabus et al.,2004b). Such methods are rather tedious, however.
To get statistically reliable data, automated analysiswas also developed, with criteria to determine whetherthe chosen set of variants in fact stems from the sameparent grain. Several ways exist, e.g. one calculation ofdeparture of MRs from well known (e.g. Burgers/Burgers) MRs [e.g. in a colonies of T40 titanium alloy(Gey and Humbert, 2003)], and iterative combination ofneighbouring product variants [e.g. in Ti–6Al–4V alloys(Glavicic et al., 2003b and 2004b)], improved by useof weighting coefficients representing the reliability ofindividual parent grain orientation data (Glavicic et al.,2003a). With such improvements, EBSD maps of theformer parent phase may even be automaticallygenerated in a robust manner and used to investigate,e.g. the solidification or forging texture of the parentphase, even after it completely disappeared duringsubsequent processing steps (e.g. in near-a IMI834(Germain et al., 2004) and in Ti–6Al–4V (Glavicicet al., 2004a) titanium alloys). A method to simplyreconstruct parent grains from EBSD data, based on arigorous algebraic analysis, that also works for fcc to bcctransformations (e.g. in steel martensite with highinternal plastic strain) has been recently developed(Cayron et al., 2006), although not yet able to calculatethe crystal orientation of the reconstructed parentgrains.
The ‘global’ approach requires little (if no) variantselection to avoid local negative values of the ODF(Humbert and Gey, 1999). In addition, the parent phaseshould not exhibit lower crystal symmetry than theproduct phase (Gey et al., 1999). In these conditions,EBSD data may be used to calculate the ODF ofproduct phase, and then (by numerical computations)that of the parent phase. Classical tools used in XRDmay readily be used. More than 2700 data points wereacquired with EBSD in Ti–6Al–4V alloy for this purpose(Humbert and Gey, 1999). A positivity criterion mayalso be introduced to allow calculations even with weakvariant selection [e.g. in metastable b Ti alloys (Geyet al., 1999; Humbert et al., 2001), tentatively inZircaloy–4 alloy (Gey et al., 2002) and in 99.85%Ti(Gey and Humbert, 2002)]. The quality of the resultsstrongly depends on the accuracy of the ODF of theproduct phase and of possible variant selection. Thismethod is much less tedious than non-automated local
approach and much less complicated than automatedlocal approach.
The calculated ODF of the parent phase may be usedin turn to calculate the ODF of the product phase byassuming no variant selection and a given OR (Humbertet al., 1994; Moustahfid et al., 1997a; Gey et al., 2002;Stanford and Bate, 2004; Glavicic et al., 2004b). Toconclude about variant selection from comparisonbetween experimental and calculated ODFs of theproduct phase, one must however bare in mind that:
(i) not all product variants were generally measuredby point analysis
(ii) analysis of 2D sections may bias the results dueto sampling effects or pseudo variant selection
(iii) local variant selection may have no effect on theresulting average texture (Stanford and Bate,2005).
Methods using only some well known texture compo-nents of the parent phase should be avoided unlessresults are experimentally confirmed by, e.g. especiallydesigned heat treatments retaining a significant part ofthe parent phase down to room temperature (Yasudaet al., 1999).
Correlation between average texture and phasetransformations
Little literature is available on average textures deter-mined by EBSD related to solid state phase transforma-tions, except for hexagonal metals and alloys (often dueto the coarse grain size of parent grains, which excludesuse of XRD) and for steels. There is a study on NiTiSMA sheets, where the final {223},110. texture resultsfrom both rolling and variant selection (Inoue et al.,2003). The effect of processing parameters on the finaltexture of product phase have also been investigated inCu–39Zn–2.6Pb (azb) brass, where appropriate extru-sion in the b phase range leads to an optimised balanceof a texture components controlling ductility andmachinability respectively (Mapelli and Venturini,2006). A strong ,110. texture influencing magneticproperties appears when Fe–9Si–13B bulk metallic glasscrystallises under a high magnetic induction (Watanabeet al., 2005a). Cyclic nitriding of 22.5Cr–5.4Ni duplexstainless steels leads to a fully austenitic case withdesirable corrosion resistant {100},001.z{110},112.
texture, while avoiding detrimental columnar grainstructure (Garzon and Tschiptschin, 2004). In diffusionbonded Ti–45Al, Ti depletion near the bond interfacelead to partial transformation of lamellar DO19 a2 phaseinto L10 c lamellae and to spheroidisation of a2 grains.While the c phase was randomly textured, new a2
particles tended to have {1010} or {0002} planes parallelto the bonding direction (Buque and Appel, 2002). Incommercially pure titanium with initial {2205},1120.
texture, the transformation from hcp a into bcc b waseventually dominated by the growth of huge allotrio-morphic b grains leading to a final {001},110. texture,even if other texture components also initially appeared(Seward et al., 2004). In acicular ferrite of steel WMdeposits, the strong ,100. solidification texture of bccd ferrite was shown by Kluken et al. (1991) in theirpioneering EBSD point analysis study to lead to threemain texture components for fcc c. As near KS ORsprevail for both d–c and c–a phase transformations, apart of the final acicular ferrite (a) microstructurerecovered the initial solidification texture. In a hot
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 95
rolled low carbon microalloyed steel, the final texture offerrite was found to be random for a low number ofpasses or a low strain per pass; it turned to typicaltexture of recrystallised ferrite with increasing strain,suggesting strain induced transformation from austeniteto ferrite, followed by deformation and recrystallisationof ferrite (Seo et al., 2000).
The average texture of phases is readily determined byEBSD in multiphase materials, without any need fordeconvolution of average pole figures, e.g. in steelscontaining some retained austenite (Field et al., 1996;Hutchinson et al., 1998; de Meyer et al., 2001;Wasilkowka et al., 2006) and in duplex stainless steels(Jura et al., 1999 and 2002). Care must be taken to assesswhether the texture of retained austenite accuratelyrepresents that of all starting austenite grains presentbefore phase transformation, and that EBSD patterns ofall phases are correctly indexed (Field et al., 1996). Thestability of retained austenite during plastic deforma-tion can also be studied for each texture componentseparately (Wasilkowka et al., 2006). Microtexture dataobtained with EBSD could here also be used inadvanced micromechanical models, to get local stressesmore accurately than by classical Taylor and Sachsapproximations.
Texture components and/or pattern quality of theparent phase provide information about its metallurgicalstate before phase transformation. By separating recrys-tallised from deformed austenite based on their wellknown texture components and by distinguishingbetween various bcc products such as primary ferrite,bainite and martensite, it was made possible to assesswhich product phase formed from a given state ofaustenite in a low carbon steel (Mesplont and DeCooman, 2003), in hot rolled dual phase steels(Waterschoot et al., 2002) and in low alloy TRIP aidedsteels (Hutchinson et al., 1998; De Meyer et al., 2001).Primary ferrite tended to form from deformed austenitewhereas bainite formed from deformed and also possiblyfrom recrystallised austenite. Local inhomogeneity inthe parent phase may strongly influence the local textureof the product phase. For instance, locally severelysheared ,110. austenite of a 0.16C–0.61Mn steeltransforms into a ,111. fibre of ferrite (Ping et al.,2005).
When comparing experimentally measured textureswith calculated ones (with assumptions about OR andvariant selection), one can find that there is nosignificant variant selection at that scale [as e.g. inspontaneous reverse transformation of warm deformedsteel (Yokota et al., 2005)]. One may even find little orno resulting texture at all [e.g. in 0.2C–0.2Si–1.3Mn–0.1Ti steel hot rolled and transformed under highmagnetic induction (Maruta and Shimotomai 2002)].Several criteria such as active slip systems or elasticstrain energy (see section on modelling for furtherdetails) have been used to include variant selection intoaverage texture calculations (Bruckner and Gottstein,2001; Humbert et al., 2002b; Moustahfid et al., 1997b;Humbert and Gey, 2003). Agreement between experi-ment and modelling has still to be improved even whentaking pseudo variant selection into account.
Solid state phase transformations: summaryWith increasing development of EBSD systems, in par-ticular for FEG-SEMs, crystallographic investigation of
solid state phase transformations is now extensively usedfrom very early growth of small, scarcely distributednuclei up to growth and coalescence of phases. At thescale of parent grains, quantitative description ofmicrotexture as ‘crystallographic’ entities, now alsoavailable for coarse grained structures, is of greatpractical significance as far as e.g. plasticity or fractureproperties are concerned. The EBSD may even replaceXRD determination of textures, especially for verycoarse grains and when several phases of close crystalstructures coexist in the investigated microstructure.Data of EBSD may be gathered and processed withroutine procedures, so that results little depend on theuser’s interpretation (except for the data ‘cleaning’procedure, which should be clearly reported). Thenumber of EBSD studies of solid state phase transfor-mations in metallic materials is exponentially increasing.With the help of EDS to facilitate phase identification,similar use of EBSD is expected for non-metallicmaterials such as ceramics and natural minerals in thefuture.
Environmentally assisted and surfacereactionsSurfaces play an increasing role in the properties ofmaterials, and in particular of functional materials.Although surface reactions are usually not referred to as‘phase transformations’, they are also considered herefor their strong similarities with phase transformationsin ‘closed’ systems, for their practical implications andbecause to the author’s knowledge, there is no recentreview of the wide use of EBSD in this field in openliterature.
Environmentally assisted surface reactionsMicrostructure and microtexture of oxide layers
Application of EBSD to the structural investigation ofoxide layers requires good spatial resolution. Oxidescales usually have a complex structure and exhibitstrong backscattered electron contrast with the under-lying metal substrate, so that cross-section examinationsrequire careful sample preparation and high perfor-mance EBSD systems with, e.g. automated backgroundacquisition during mapping.
A significant practical application of EBSD in thisfield is the investigation of oxide scales on steel products,whose structure and mechanical properties are of primeimportance for hot rolling, cold rolling and wiredrawing. Both FeO1–x (wustite), Fe3O4 (magnetite) andFe2O3 (hematite) may be distinguished with EBSD,although there is usually little Fe2O3 (Kim and Szpunar,2001) or the Fe2O3 layer is too thin to be detected withEBSD (Burke and Higginson, 2000; Birosca andHigginson, 2003). Interstitial free and low carbon steelsare of primary concern. According to oxidation condi-tions (temperature, atmosphere and cycling conditions)and steel composition, the internal wustite layer may beeither columnar (Burke and Higginson, 2000; Kim andSzpunar, 2001 and 2002; West et al., 2005) or equiaxed(Birosca and Higginson, 2003; Birosca et al., 2004). Astrong ,100. texture may be encountered in wustitewhatever that of the steel substrate (Kim and Szpunar,2001 and 2002; Higginson et al., 2002). A cube–cube ORbetween wustite and magnetite may also prevail,possibly due to the defect structure of wustite (West
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
96 International Materials Reviews 2007 VOL 52 NO 2
et al., 2005), but is not always found (Birosca andHigginson, 2003). Silicon significantly affects the struc-ture and texture of the oxide layers (Kim and Szpunar,2002; Higginson et al., 2002). The oxide texture may becorrelated with porosity and thus mechanical propertiesof oxide layers (Higginson et al., 2002). Scales producedin ‘industrial’ conditions have also been characterisedwith EBSD (Burke and Higginson, 2000). The limits ofsuch analysis are spatial resolution (for very thin scalesto be observed in cross-sections) and statistical reliabilityof data. However, sample preparation is much lesstedious than for TEM cross-sections and usually doesnot need to use advanced techniques such as focused ionbeam (FIB) milling.
Oxide scales have also been studied in cross-sectionsof T40 and Ti–3Al–2.6V titanium alloys, where rutileTiO2 grows faster in Ar rich atmospheres than in N2 richatmospheres and a tendency to (0002)Ti//(100)TiO2 wasobserved (Lenarduzzi et al., 2002). Thermal barriercoatings on nickel base superalloys have also beenstudied with EBSD, for various coating processes. Thealumina scale may, or may not be strongly ,0001.
textured. This controls internal stresses and finalproperties of the thermal barrier coatings (Karadgeet al., 2006). The development of internal stress andstrain fields was investigated with EBSD after oxidationof buried AlGaAs layers; residual stresses and plasticzone geometry were calculated by a finite element
method and compared to the distribution of patternquality (Keller et al., 2004).
Another way to investigate the microtexture of layersis to carry out EBSD analysis of oxidised surfacesdirectly (Fig. 15). This provides data over a wide area,although only at the external surface. Oxidation ofrecrystallised {100},001. (cube) textured nickel lead totwo types of areas:
(i) rough areas with coarse cube or {111} texturedgrains and smaller, possibly faceted ‘intergranu-lar’ grains
(ii) smooth regions having the desired epitaxial cubeor rotated cube texture (Woodcock et al., 2002a,2002b and 2004).
Role of interfaces and free surfaces
As environmentally assisted mechanisms are closelyrelated to diffusion, interfaces and free surfaces stronglyinfluence morphology and growth kinetics of products.As there is already review literature on EBSD studies ofintergranular cracking (Gourgues, 2002), only corrosionis addressed here (Table 22). Most results were obtainedby analysis of the corroded surface and of parallel planesprepared by serial sectioning, sometimes confirmed bycross-section analysis.
The crystal orientation of free surface planes, which isreadily determined with EBSD, also influences the layergrowth kinetics. The relevant parameter is generally the
(a) (b) (c)
a prior EBSD mapping of surface to be oxidised; b exposure to oxidising environment; c characterisation of oxidisedsurface (e.g. with near field or interferometry techniques)
15 Principle of EBSD investigation of oxidation
Table 22 Some EBSD studies of intergranular corrosion
Material Corrosion conditions EBSD results Ref.
Ni–39–40Fe (at.-%) 1000uC for 5 h in O2
atmosphereRandom GBs are sensitive and LABs are not sensitiveto corrosion; the sensitivity of coincident site lattice GBsincreases with oxygen pressure and with departure fromS3, S11, S19 and S27 coincidence
Yamaura et al.,1999 and 2000
Ni–39 at.-%Fe 800–927uC for 18–24 hin air
Sensitivity increases with applied stress; 65% of randomGBs in the sample is enough to oxide over the full width
Yamaura et al.,2003
Low carbon ship steels Seawater environment Not all GBs are corroded, some can only be imaged byEBSD
Katrakova andMucklich, 2000
AA6061 aluminium alloy Intergranular attack inseawater environment
LABs are immune; intergranular attack stops when thereis a loss in connectivity of sensitive HABs
Minoda andYoshida, 2002
Sensitised AISI 316LNstainless steel
Intergranular attack inoxalic acid
LABs and twin boundaries are not sensitive; misorientationangle of sensitive GBs spreads between 30 and 55u; suchboundaries form a continuous network
Kunıkova et al.,2004
Nickel base alloy 718 900uC for 30 h in air Twin boundaries are immune; random GBs are oxidised Yang et al.,2005
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 97
angle between the local free surface and low-indexcrystal planes such as {0001} planes in oxidised Grade 1titanium (Konig and Davepon, 2001) and {100} planesin nickel base alloy 690 subjected to plasma assistednitriding (He et al., 2003; Czerwiec et al., 2003).Transmission electron microscopy results are difficultto obtain, whereas EBSD may be conducted before theenvironmentally assisted reaction, without affecting thesample geometry. This has been done for Co51.6Ga48.3
oxidised into Ga2O3, for which the oxide growth kineticson {211} substrate planes was about five times higherthan that on {100} planes (Koops et al., 2002). In nickelbase alloys, atomic force microscopy of corrodedsurfaces (characterised by EBSD before corrosion)allowed quantitative assessment of the dissolution rateas a function of the local surface orientation (Gray et al.,2006; Schuh et al., 2003) and crystallographic investiga-tion of fine scaled surface features of individualcorroded grains (Schuh et al., 2003).
Strong interactions between substrate GBs and freesurfaces have been evidenced thanks to EBSD forwustite reduction into iron in CO/CO2 atmospheres bythe following reactions
CO adð Þz2 h:zV ’’
FezOXO?CO2 adð Þ (1)
FeXFe?Fe0z2 h
:zV ’’
Fe (2)
Reaction (1) leads to local depletion in oxygen and thento nucleation of metallic iron by reaction (2). It alsoleads to surface rearrangement, in particular near GBs(Bahgat et al., 2004a). Surface rearrangement leads tolocal roughness depending on the individual wustitegrain orientation (Bahgat et al., 2004b and 2005). Evenwhen still not visible in the SEM, iron nuclei can beclearly identified by EBSD, since EBSD is sensitive tothe very surface of the specimens (Sasaki et al., 2005b).Iron particles nucleate at surface ledges and their sizeand number strongly depend on the local wustite grainorientation (Bahgat et al., 2004a and 2004b). Thegrowth of iron particles also depends on the localrearranged surface (Sasaki et al., 2005b). Such studiesmake full use of the non-destructive, highly resolved inthickness EBSD technique.
Other results on environmentally assisted reactions
A variety of reactions have been explored with EBSDfor either phase identification or determination of ORand microtexture (Table 23). Another interesting exam-ple is given by in situ reaction synthesis of alumina fromreaction between molten aluminium and surroundingSiO2 quartz tube (Murthy et al., 2005). Al2O3 and Alwere found as fully interconnected colonies with onlytwin boundaries between Al2O3 grains, which should bebeneficial to mechanical properties of the final product ifcomplete synthesis could be obtained.
Thin film depositionEBSD has been used at all stages of process develop-ment for thin film deposition, from feasibility demon-stration to quality control of the final product.
Process feasibility
Electron backscatter diffraction has been used to findout processing conditions yielding the desired phase,such as amorphous versus crystalline silicon on singlecrystal silicon substrate (Gao et al., 2000), crystalline
silicon by metal assisted nucleation in amorphous Si:Hlayer (Chang et al., 2004) and silicide mediated versuslaser assisted crystallisation of silicon on glass (Kimet al., 2004). Quantitative determination of the amountof given phases or polytypes may also be given by EBSDanalysis of deposited films, such as SiC deposition onto4H and 6H SiC substrates (Chaussende et al., 2004;Latu-Romain et al., 2005). Once the right phase isobtained, the next point is to control its crystal qualityand orientation. Crystal quality may be assessed byinspection of GBs or by looking at EBSD patternquality (Kim et al., 2004). Controlling the film texture(e.g. through epitaxial growth combined with variantselection) can be more difficult than controlling thenature of polytype (Chaussende et al., 2004). Exceptwhen deposited films were thick enough [e.g. TiN innickel base superalloys (Jeong et al., 2002b)], only freesurfaces of films were generally investigated, making fulluse of the sensitivity of EBSD to extreme surface layersonly.
Some results given by EBSD analysis about thefeasibility of obtaining suitably oriented films on varioussubstrates are illustrated in Table 24. Extensive mappingwas not always necessary, simple eye observation ofchanges in EBSD patterns by rastering the electronbeam over the entire sample surface being preferred inmany cases.
Orientation relationships between layers and substrate
The OR between substrate and deposited layers has beeninvestigated with EBSD in a variety of cases, mainly bysurface analysis of films (Table 25). In some cases,however, sampling or phase size effects may causemissing of variants that were indeed found by XRD(Cain and Lange, 1994).
Layer structure and resulting texture
Spatial information provided by EBSD has been usedto determine the grain size, variant morphology andclustering and even grain morphology on polished cross-sections. Data are now available for various materialsand deposition processes (Table 26). The resultingtexture is usually analysed with XRD but has also beeninvestigated with EBSD (Table 27).
Information on local roughness has been used to linkthin film morphology to EBSD determined crystal struc-ture and orientation of substrate and film. One may cite,e.g. copper electrodeposits (Cho and Szpunar, 2002),dendritic silicon splats (Nagashio and Kuribayashi,2005), CeO2 buffer layers (Van Driessche et al., 2003),NiO grown on Ni substrate (Woodcock et al., 2004), SiC(3C) grown on SiC (6H) (Latu-Romain et al., 2005), andsputter deposited Nb, Cu, Co and Permalloy epitaxiallayers (Loloee et al., 2001).
Surface reactions: a summaryEssential features of the EBSD technique in the field ofsurface reactions are as follows:
(i) its sensitivity to extreme surface layers allowssurface characterisation of as deposited extre-mely thin films (down to y20 nm in thickness)
(ii) its non-destructive character is essential to deter-mine ORs between substrate and layers forpolycrystalline substrates
(iii) the possibility is given to scan over large areas ata fine scale.
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
98 International Materials Reviews 2007 VOL 52 NO 2
Tab
le23
Ele
ctr
on
backscatt
er
dif
fracti
on
stu
die
so
fen
vir
on
men
tall
yassis
ted
su
rface
reacti
on
s
Mate
rial
En
vir
on
men
tS
urf
ace
reacti
on
EB
SD
resu
lts
Ref.
Pure
Mg
0. 0
1M
NaC
lz10
–4M
Na
2C
r 2O
7
aq
ueous
solu
tion
Fila
menta
rycorr
osio
nC
orr
osio
nin
itia
tes
only
on
{0001}
oriente
dg
rain
s.
Surf
ace
pro
pag
ation
alo
ng
,1120
.and
,1010
.
directions
Schm
utz
et
al.,
2003
Ti–
8at.
-%M
oN
itrid
ing
environm
ent
Inte
rnalTiN
(hcp
a)
nitrid
ing
Burg
ers
OR
with
matr
ix,
1–3
variants
per
bcc
bg
rain
of
the
sub
str
ate
Guill
ou
et
al.,
2004
Ti–
16
at.
-%M
oN
itrid
ing
environm
ent
Inte
rnalTiN
1–x
(d,
‘fcc’)
nitrid
ing
Near-
KS
OR
,M
Rs
betw
een
part
icle
sare
oft
en
60u,
111
.
Guill
ou
et
al.,
2004
Pure
iron
and
two
low
carb
on
low
allo
yste
els
Borid
ing
with
or
without
carb
urisin
gIn
tern
alfo
rmation
of
tetr
ag
onalFeB
2
part
icle
sS
ing
lecry
sta
lline
Fe
2B
part
icle
sin
borid
ed
Fe;
poly
cry
sta
lline
Fe
2B
part
icle
sin
borid
edz
carb
urised
ste
els
Kulk
aet
al.,
2006
Dup
lex
sta
inle
ss
ste
els
Nitrid
ing
at
1200uC
Surf
ace
dis
solu
tion
of
bcc
dp
hase
Identification
of
the
cp
hase
and
of
dis
ap
peara
nce
of
dp
hase
Pad
ilha
et
al.,
1999
Fe–4
at.
-%N
i–(2
%Tior
3%
Cr)
Nitrid
ing
environm
ent
(thin
sheet)
Part
ially
revers
ible
bulk
transfo
rmation
into
nitrid
es
and
resultin
gg
rain
refinem
ent
Phase
identification
(EB
SD
),g
rain
siz
eand
morp
holo
gy
(TE
M)
Chezan
et
al.,
2004
Fe–18
wt-
%C
rP
lasm
anitrid
ing
Nitrid
ep
recip
itation
and
cellu
lar
pre
cip
itation
Cry
sta
llog
rap
hic
continuity
betw
een
sub
str
ate
gra
ins
and
gra
ins
of
the
nitrid
ed
layer
Miy
am
oto
et
al.,
2006
Fe–28M
n–6S
i–5
wt-
%C
r900uC
for
1h
ina
vacuum
of
10
–2
Pa
Evap
ora
tion
of
Mn
causin
gexte
rnal
fcc
cto
bcc
atr
ansfo
rmation
Phase
identification
of
aFukaiet
al.,
2005
Nib
ase
allo
y22
NaC
lzH
Claq
ueous
solu
tions
Com
petition
betw
een
passiv
ation
and
meta
ld
issolu
tion
3M
HC
l:d
issolu
tion
accord
ing
toth
elo
calnorm
alto
free
surf
aces:
incre
asin
gra
tes
for
{111},
{100},
{110}
1M
HC
l:re
sid
ualoxid
ation
com
pete
sw
ith
dis
solu
tion;
resultin
gra
te{1
11},
{110},
{100}
(quantita
tive
descrip
tion
inin
vers
ep
ole
fig
ure
s)
Gra
yet
al.,
2006
Nib
ase
allo
y600
0. 1
NH
Cld
rop
let
Dis
solu
tion
and
oxid
ation
,100
.b
and
son
corr
od
ed
{110}
surf
aces.
Dis
solu
tion
rate
:{1
11},
{110},
{100}
(quantita
tive
descrip
tion
inin
vers
ep
ole
fig
ure
s)
Schuh
et
al.,
2003
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 99
Tab
le24
So
me
EB
SD
resu
lts
on
qu
ali
tyo
fth
infi
lmd
ep
osit
sfo
rfe
asib
ilit
yd
em
on
str
ati
on
Su
bstr
ate
Film
Dep
osit
ion
pro
cess
Resu
lts
Ref.
(100)
ZrO
2–9. 5
%Y
2O
3(c
ub
ic)
ZrO
2A
queous
pre
curs
orz
cert
ain
annealin
gcond
itio
ns
Sam
eE
BS
Dp
attern
over
the
entire
sam
ple
surf
ace
Mill
er
et
al.,
1993
(100)
Cu
(fcc)
Cu
and
Co–N
im
ultila
yers
(0. 1
5nm
per
layer)
Pote
ntiosta
tic
ele
ctr
od
ep
ositio
nE
pitaxy
(cub
e–cub
eO
R)
Alp
er
et
al.,
1993
Rolle
dand
polis
hed
Ni–
Wallo
yC
eO
2S
ol–
geld
ipcoating
Cub
e–cub
eorienta
tion
with
poly
cry
sta
lline
sub
str
ate
Van
Driessche
et
al.,
2003
(0001)
GaN
,N
ith
en
Au
Meta
llisation
Cry
sta
lq
ualit
yof
both
layers
invarious
cond
itio
ns;
suitab
leep
itaxy
ifN
i,th
en
Au
are
dep
osited
Davyd
ov
et
al.,
2004
(1120)
Al 2
O3
(sap
phire)
Nb
DC
mag
netr
on
sp
utt
ering
at
750uC
Clo
ser
than
,1u
toep
itaxy
Lolo
ee
et
al.,
2001
Tl 2
Ba
2C
a2C
u3O
xLaA
lO3
DC
mag
netr
on
sp
utt
ering
Cub
e–cub
eep
itaxy
Bra
mle
yet
al.,
1998
(001)
and
(111)
Mg
O*
In2O
3P
uls
ed
laser
dep
ositio
nz
anneal
Ep
itaxy,
OR
bein
g1. 7
¡0. 5u
,111
.*
Johnson
et
al.,
1999a
(100)
SrT
iO3
(Pb
0. 5
2Z
r 0. 4
8)T
iO3
Puls
ed
laser
dep
ositio
nC
ub
e–cub
eep
itaxy{
Ham
ed
iet
al.,
1998
Siof
various
orienta
tions
Si
Ele
ctr
on
cyclo
tron
resonance
chem
icalvap
our
dep
ositio
n(1
00)
Si:
ep
itaxy;
(210)
Si:
ep
itaxy
with
defe
cts
;(3
11),
(111),
(221):
no
ep
itaxy
Rau
et
al.,
2004
(0001)
Al 2
O3
(sap
phire)
Al 2
O3
Solid
sta
teconvers
ion
of
Al
All
convert
ed
Al 2
O3
part
icle
sare
of
sam
ecry
sta
lorienta
tion
as
sub
str
ate
(nanohete
roep
itaxy)
Park
et
al.,
2005
Anod
ised
poly
cry
sta
lline
Ti
Pb
O2
Ele
ctr
od
ep
ositio
nFeasib
ililit
yonly
on
{0001}
Tig
rain
s,
too
difficult
oth
erw
ise
Devill
iers
et
al.,
2004
*In
vestig
ation
incro
ss-s
ection.
{ Surf
ace
investig
ation
by
ele
ctr
on
channelli
ng
patt
ern
analy
sis
(clo
se
toth
eE
BS
Dte
chniq
ue).
Tab
le25
Ele
ctr
on
backscatt
er
dif
fracti
on
investi
gati
on
of
ori
en
tati
on
rela
tio
nsh
ips
betw
een
thin
dep
osit
ed
layers
an
dsu
bstr
ate
Su
bstr
ate
Film
Dep
osit
ion
pro
cess
Resu
lts
Ref.
(1120)
Al 2
O3
(sap
phire)
Nb
DC
mag
netr
on
sp
utt
ering
at
750uC
(110) N
b//
(1120) A
l 2O
3and
[111] N
b//
[0001]
Al 2
O3
by
,1u
Lolo
ee
et
al.,
2001
Nb
(of
pre
ced
ing
row
)(b
cc)
Cu
(fcc)
DC
mag
netr
on
sp
utt
ering
at
350uC
or
(150uC
zannealat
350uC
)N
WO
R,
with
two
variants
Lolo
ee
et
al.,
2001
(0001)
Al 2
O3
(sap
phire)
rota
ted
by
10u
tow
ard
(1010)
GaN
(30
nm
)C
hem
icalvap
our
dep
ositio
n(1
010) G
aN
//(1
210) A
l 2O
3i.e.
90u
aro
und
,0001
.Tra
ger-
Cow
an
et
al.,
2001
(0001)
and
(1120)
Al 2
O3
(sap
phire)
ZrO
2–10
mol.-%
Y2O
3A
queous
pre
curs
or
(0001):
thre
evariants
(001) Z
rO2
//(0
001) A
l 2O
3and
,100
.Z
rO2
//,
1210
.A
l 2O
3
with
one
dom
inating
variant;
(1120):
four
variants
with
various
OR
s,
som
eoth
ers
dete
cte
db
yX
RD
only
Cain
and
Lang
e,
1994;
Cain
et
al.,
1995
Poly
cry
sta
lline
Cu
Cu
dro
ple
tsA
nnealat
1050uC
inH
2/H
eatm
osp
here
There
are
pre
ferr
ed
initia
lO
Rs
but
with
limited
effect
on
sp
read
ing
.E
pitaxy
at
the
end
of
sp
read
ing
(GB
mig
ration).
Mis
sia
en
et
al.,
2005
GaN
Ni
Ele
ctr
on
beam
dep
ositio
n[1
10] N
i//
[1120] G
aN
or
[1210] G
aN
Davyd
ov
et
al.,
2004
(110)
TiO
215
nm
siz
ed
Au
part
icle
sIn
situ
meta
llisation
and
anneal
at
527uC
80%
of
101
analy
sed
part
icle
shave
(111)
,10u
from
(110) T
iO2;
two
variants
are
found
with
(111) A
u//
(110) T
iO2
and
(I)
[110] A
u//
[001] T
iO2
and
(II)
[110] A
u//
[001] T
iO2
Cosand
ey,
1997
(110)
TiO
212
nm
siz
ed
Au
part
icle
sIn
situ
meta
llisation
at
20uC
(zanneal)
or
502uC
20uC
:(1
11) A
u//
(110) T
iO2b
y3u;
two
variants
are
found
(I)
[110] A
u//
[001] T
iO2
and
(II)
[110] A
u//
[001] T
iO2
502uC
:(1
12) A
u//
(110)
TiO
2b
y2u
and
[110] A
u//
[001] T
iO2
or
[110] A
u//
[001] T
iO2
or
twin
rela
ted
with
these
Cosand
ey
et
al.,
2001
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
100 International Materials Reviews 2007 VOL 52 NO 2
Tab
le26
Usin
gE
BS
Dto
dete
rmin
eth
estr
uctu
reo
fd
ep
osit
ed
layers*
Su
bstr
ate
Film
Dep
osit
ion
pro
cessz
ob
serv
ati
on
co
nd
itio
ns
Resu
lts
Ref.
(001)
Mg
OFe
2O
3P
uls
ed
laser
dep
ositio
nz
(IP
)Four
self-a
ccom
mod
ating
variants
gro
up
ed
into
recta
ng
ula
rd
om
ain
s(w
ith
deta
iled
mic
rote
xtu
re)z
possib
lein
tern
altw
innin
gJohnson
et
al.,
1999b
(100)
Sicoate
dw
ith
Ti,
then
Ni,
then
(111)
Pd
Sn–3. 5
%A
g(b
-Sn
phase)
Sta
ck
ele
ctr
op
latingz
(CS
)(2
20)
textu
re,
avera
ge
gra
insiz
ey
21mm
zevolu
tion
during
annealin
gE
zaw
aet
al.,
2004
(?)
YB
a2C
u3O
7–
dD
Csp
uttering
at
780uC
z(I
P)
Gra
insiz
ed
istr
ibution
(y1mm
2),
desirab
le[0
01]
textu
reob
tain
ed
at
98. 5
%Fairhurs
tet
al.,
2000
Pure
poly
cry
sta
lline
Cu
of
various
textu
res
Cu
(50mm
)E
lectr
od
ep
ositio
nz
(CS
)C
oars
eep
itaxia
lcolu
mnar
gra
ins,
com
petitive
gro
wth
ztw
innin
gC
ho
and
Szp
unar,
2002
,111
.S
icoate
dw
ith
60
nm
Cu
Cu
(2. 5
mm
)E
lectr
od
ep
ositio
nz
(IP
)G
rain
siz
ey
1. 5
mm
much
hig
her
than
siz
eof
ind
ivid
ualsp
here
sR
ead
et
al.,
2004
Nic
kelb
ase
sup
era
lloy
TiN
(6mm
)P
hysic
alvap
our
dep
ositio
nz
(CS
)S
mall
(0. 1
5–1. 5
mm
)colu
mnar
gra
ins,
gra
insiz
ed
istr
ibution
Jeong
et
al.,
2002b
Mo
Dia
mond
Chem
icalvap
our
dep
ositio
nin
recycle
dg
as
at
830uC
z(I
P,
CS
)C
om
petitive
gro
wth
and
twin
nin
gand
associa
ted
textu
reevolu
tion
inth
ickness
Mao
et
al.,
2005
(1120)A
l 2O
3(s
ap
phire)
coate
dw
ith
(110)
Nb
Perm
allo
y(N
i–16Fe),
then
Cu
DC
mag
netr
on
sp
utteringz
(IP
)C
uand
Perm
allo
y:
two
variants
with
180u,
111
.M
R,
4–5mm
insiz
eLolo
ee
et
al.,
2001
Poly
cry
sta
lline
Ni 3
Al
Dia
mond
Pla
sm
aassis
ted
chem
icalvap
our
dep
ositio
nw
ith
positiv
eb
ias
enhanced
nucle
ationz
(IP
)N
ucle
ation
density
dep
end
son
sub
str
ate
gra
inorienta
tion
Chen
and
Chang
,2005
(001)
SrT
iO3
coate
dw
ith
IrD
iam
ond
Pla
sm
aassis
ted
chem
icalvap
our
dep
ositio
nw
ith
bia
senhanced
nucle
ationz
(IP
)N
ucle
ation
str
uctu
res:
sam
eE
BS
Dp
attern
as
Irb
ut
are
infa
ct
futu
red
iam
ond
dom
ain
s;
the
reaction
isauto
cata
lytic
Schre
ck
et
al.,
2003
,111
.S
i nD
iam
ond
Pla
sm
aassis
ted
chem
icalvap
our
dep
ositio
nw
ith
bia
senhanced
nucle
ationz
(serial
sections
para
llelto
IP)
Colu
mnar,
mostly
ep
itaxia
lg
rain
sC
hen
and
Rud
olp
h,
2003
No
sub
str
ate
SiC
(3C
)Flu
idis
ed
bed
dep
ositio
nfr
om
CH
3S
iCl 3
(CS
)C
olu
mnar
gra
ins
from
the
eq
uia
xed
chill
zone;
inte
rnaltw
innin
g;
alm
ost
no
avera
ge
textu
reH
ela
ryet
al.,
2006
Ti
Nanostr
uctu
red
Co–20
at.
-%N
iE
lectr
od
ep
ositio
n(C
S,
IP)
5%
fcc
(Co)
phase
with
OR
tohcp
matr
ix;
gra
insiz
e,
GB
chara
cte
risation;
cry
sta
lorienta
tion
gra
die
nt
incolu
mnar
gra
ins
Basto
set
al.,
2006
AIS
I316
sta
inle
ss
ste
el
Pd
Ele
ctr
od
ep
ositio
nz
(CS
)G
rain
siz
ed
istr
ibution
(140
nm
inavera
ge),
gro
wth
from
sub
str
ate
gra
ins
far
from
GB
sB
era
et
al.,
2004
*IP
:in
pla
ne
ob
serv
ations;
CS
:ob
serv
ation
of
cro
ss-s
ections.
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 101
For in plane analysis, a balance must be found betweenspatial resolution and sampling effects and the rough-ness of the analysed surface must be low. Cross-sectionanalysis may require specific sample preparation (e.g. byFIB milling techniques) to avoid edge effects.
Discussion: coupling EBSD with otherinvestigation methodsAs EBSD mapping has become in most cases a routinetechnique, it has widely been used together with otherexperimental characterisation techniques for multiscaleinvestigation of materials. The case of EBSD/EDXcoupling for phase identification has already beenaddressed in the present paper. On the other hand,information on coupled crystal orientation and mor-phology is used in advanced models, which are now evenable to take into account individual or clustered crystalsto optimise process conditions or product properties.
Coupling EBSD with other experimentaltechniquesValidation of experimental results
In mineralogy samples, the orientation of individualcrystals is generally investigated with the light micro-scope. Advantages and drawbacks of EBSD and lightmicroscopy respectively, as well as coupled use of bothhave already been reviewed (Trimby and Prior, 1999).The main advantages of EBSD are its ability to analysesmall grains, optically ‘isotropic’ crystals and transpar-ent materials. Grain boundaries are more quantitativelycharacterised than by light microscopy (which is notsensitive to LABs) or backscattered electron imaging(Trimby and Prior, 1999). However, a robust reflectiontable must be constructed for each analysed mineral(Mauler et al., 1998).
The EBSD technique has also been used together withFIB imaging. In Al–0.5Cu, the ion channelling con-trast was shown with EBSD to rely on the angularmisorientation between the local normal to the samplesurface and the nearest ,111. crystal direction (Barret al., 1992).
Coupled analysis with EBSD and near field micro-scopy has been performed for less than a decade forphase identification [e.g. to distinguish ferrite fromaustenite in multiphase steels (Wendrock et al., 2001;Ros-Yanez et al., 2001 and 2002)], and to investigate GBgrooving [e.g. shallow thermal grooves at LABs orcoincident site lattice boundaries of polycrystalline MgO(Farrer et al., 2000)]. The high spatial resolution ofatomic force microscopy has been used to checkaccuracy of EBSD determination of GB traces in TiO2
polycrystals (Pang and Wynblatt, 2006), and to checkthat sapphire crystals obtained by solid state conversionwere of sufficient height (here y100 nm) for theirorientation to be determined with EBSD independentlyof that of the sapphire substrate (Park et al., 2005).
Multiscale analysis
Although EBSD makes full use of the large magni-fication range of the SEM, there is still a need for finerscale analysis and for analysis of higher amounts ofmaterial to reduce sampling effects. The coupled use ofvarious techniques for that purpose is discussed in thissection.T
ab
le27
Usin
gE
BS
Dto
dete
rmin
eavera
ge
textu
reo
fd
ep
osit
ed
layers
Su
bstr
ate
Film
Dep
osit
ion
pro
cess
Resu
lts
Ref.
Poly
(meth
ylm
eth
acry
late
)coate
dw
ith
Cu
Ni
Ele
ctr
od
ep
ositio
nin
sulfam
ate
(S)
or
Watts
bath
(W)
(S):
,100
.colu
mnar
zone;
(W):
fine,
poorly
ind
exed
mic
rostr
uctu
re,
weaker
,110
.
Buchheit
et
al.,
2002
Alcoate
dw
ith
Zn
then
with
Cu
Ni
Ele
ctr
od
ep
ositio
n,
110
.colu
mnar,
many
low
mis
orienta
tions
within
colu
mns
Arn
ould
et
al.,
2003
Ti
Co–20
at.
-%N
iE
lectr
od
ep
ositio
nS
trong
{1120} a
and
{220} b
textu
reB
asto
set
al.,
2006
Mg
OTl 2
Ba
2C
a2C
u3O
xD
Cm
ag
netr
on
sp
uttering
(001)
fib
reand
many
HA
Bs:
poor
film
pro
pert
ies
Bra
mle
yet
al.,
1998
Sicoate
dw
ith
SiO
2th
en
with
Tith
en
with
Pt
then
with
La(N
O3) 3
zp
ossib
leb
uff
er
layers
(Pb
,Zr)
TiO
3S
ol–
gelcoating
Str
ong
(100)
textu
re,
no
textu
rein
pla
ne
Choiet
al.,
2004
and
2005
Mo–9S
i–18
at.
-%B
MoS
i 2z
transfo
rmation
into
Mo
5S
i 3d
uring
annealin
g
Pack
cem
enta
tionz
annealin
gM
oS
i 2:
[001]
at
exte
rnalsurf
ace;
Mo
5S
i 3:
[001]
colu
mnar
Ito
et
al.,
2003
Oxid
ised
(100)
Sicoate
dw
ith
100
nm
Cu
then
with
282
nm
Al
Al 2
Cu
Ele
ctr
on
beam
evap
ora
tion
Cellu
lar
gro
wth
with
pre
fere
ntial(1
10)
textu
reS
on
et
al.,
2001
No
sub
str
ate
TiO
2(r
utile
)Tap
ecastingz
sin
tering
TiO
2:
(001)
textu
re,
[001]
gro
wth
direction
Cosand
ey
et
al.,
1999
No
sub
str
ate
Al 2
O3
Tap
ecastingz
sin
tering
(0001)
para
llelto
both
film
and
pla
ne
of
pla
tes
Mark
ond
eya
Rajet
al.,
1999
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
102 International Materials Reviews 2007 VOL 52 NO 2
Additional information provided by coupling EBSD withnear field microscopy
Near field microscopy gives access to local 3D surfacemorphology and material properties at a very fine scale.Shreck et al. (2003) were able to detect diamond nucleion iridium substrate by Kelvin probe microscopy,although the local crystal structure was still determinedby EBSD to be that of iridium. Magnetic forcemicroscopy readily detected magnetisation domains atGBs of AISI 304 austenitic stainless steel, which wereconfirmed by fine scale EBSD mapping to be martensiteparticles (Takaya et al., 2004). Local surface modifica-tion after thermal grooving [e.g. in SrTiO3 (Sano et al.,2003a)] gives information about free surface and GBenergy.
Local surface relief produced on polished surfaces byphase transformation (e.g. from bcc b to orthorhombica’’ in Ti–9 wt-%Mo (Guo and Enomoto, 2006), fromaustenite into Widmanstatten ferrite (here characterisedwith a specific procedure by light microscopy) (Hall andAaronson, 1994) or into bainite (Ohmori, 2002) in steel)suggests surface reconstruction or previous twinning ofthe parent phase near single crystal, tent shaped plates(i.e. a diffusional mechanism) and no diffusion nearplates showing an invariant plane strain morphology.More generally, the macroscopic strain induced bydisplacive solid state phase transformations at polishedsurfaces can be determined in 3D using near fieldmicroscopy, whereas local crystallography in the samearea is determined with EBSD. Such precious data areessential for validation of phenomenological theoriesof martensite crystallography (PTMCs) (see Modellingsection).
The measurement of local properties coupled withEBSD determination of the local crystal structure andorientation has been recently reviewed in the frame-work of phase diagram determination (Zhao, 2006).Nanoindentation on individual phases identified withEBSD was performed in Pd–Rh–Pt diffusion multiples(Zhao et al., 2002), on Ag3Sn and Cu6Sn5 intermetallicphases in Sn–Ag–Cu solder interconnects (Li et al.,2005b) and in various layers formed on a nickel basesuperalloy coated with NiAl by pack cementation(Wollmer et al., 2003). The anisotropy of propertiesmay be quantified thanks to information about localcrystal orientation [e.g. in the NiAl phase of NiAl–Moeutectic (Bei and George, 2005)]. There are cases forwhich a simple rule of mixtures between the variousphases worked well for the investigated microstructure[e.g. hardness and elastic modulus of a V–V3Si eutectic(Bei et al., 2004)]. In other instances, no agreement couldbe found, e.g. for elastic modulus and for hardness(Zhao et al., 2002). Correlation between nanohardnessand local loss of EBSD pattern quality is still in progress(Wu et al., 2005).
EBSD versus TEM investigations
A huge amount of studies use both TEM and SEM-EBSD investigations of the same phenomena; auto-mated indexing of Kikuchi bands is now possible in theTEM, allowing orientation imaging at very fine scales.This section focuses on complementary use of SEM-EBSD and TEM techniques. Transmission electronmicroscopy investigations are needed for fine scalecharacterisation of deformation structures, such as
dislocation clustering into LABs that are observed withEBSD but do not affect the martensitic transformation[e.g. in Cu–Al–Ni SMA (Rodriguez, 2004)], and if thereare local deformation bands in the parent phase, inwhich particular variant selection may then occur [e.g.in Cu–42 wt-%Zn (Sakata et al., 2000)]. Interactionsbetween the local texture and very fine scale precipita-tion also require TEM observations [e.g. precipitationin Zircaloy–4 (Loge et al., 2000)]. The effect of GBmisorientation on discontinuous precipitation in Al–2.8Mg–1 at.-%Ga was determined with EBSD, whereasthe ORs between phases had to be determined withTEM (Hirth and Gottstein, 1998). In precipitation aidedmatrix phase transformations, the OR between matrixphases may be determined with EBSD while the ORswith fine nucleating agents are still studied in the TEM[e.g. ferritic steels (Furuhara et al., 2003)].
Fine scale microstructural features (such as thosedifferentiating ferrite from bainite in low alloy TRIPsteels) were first studied by TEM, giving qualitative orquantitative criteria for further (and easier) investigationby SEM-EBSD. In this particular case (Fig. 16),orientation gradient and internal LABs were observedin bainite, but not in ferrite (Zaefferer et al., 2004). Theinternal structure of c phases in c-TiAl gave comple-mentary information to EBSD determined ORs betweenlamellar and Widmanstatten c phases in this alloy (Deyet al., 2005). The gradual misorientation determinedby EBSD within lenticular a’ martensite of Fe–32.85 wt-%Ni (Fig. 17) was shown by TEM to be linkedto a change in the internal structure of martensite (andlocal accommodation of transformation induced latticestrains) from dislocation network to microtwinning(Shibata et al., 2005). Internal microtwinning in self-accommodating martensite of Cu–12.55Al–4.84 wt-%Ni(Chen et al., 2000) and in Cu–7.3Al–8.5 wt-%Mn (Wanget al., 2002) was detected in the TEM but not by SEM-EBSD, so that only a part of the self-accommodatingvariants could be identified in each martensite unit byEBSD. TEM based techniques remain necessary if veryfine scale features or dislocation structures have to beinvestigated.
Other diffraction techniques
X-ray diffraction is a reference technique for texturedetermination, so that the comparison of resultsobtained by XRD and EBSD respectively, sincepioneering studies of e.g. Schwarzer and Weiland ondual phase steels (Schwarzer and Weiland, 1988), hasalready been reviewed (Dingley and Randle, 1992). Inaddition to providing spatial information, EBSD is wellsuited for coarse grained materials. Lattice parametersare better determined with XRD, and a large area can beanalysed at a time. The probability to miss particularvariants or phases is much lower with XRD than withEBSD [e.g. in cubic ZrO2 deposited onto sapphire (Cainand Lange, 1994)]. For moderately coarse microstruc-tures or for very low amounts of phases, neutrondiffraction coupled with EBSD is a well suited tool. Thepioneering study of Grant et al. (1986) focused ondirectionally solidified copper, where optical microscopyrevealed a columnar grain structure, neutron diffractionshowed a ,100. fibre texture that was confirmed bylocal electron channelling pattern analysis. In anaustenitic stainless steel WM deposit, Bouche et al.(2000) showed by both neutron diffraction (providing
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 103
statistics over y1 cm3 of material), point EBSD analysisof submicron sized bcc ferrite and local TEM observa-tions that the OR between austenite and ferrite was notonly the classical near KS OR, but could be much closerto a cube–cube OR, at least for ferrite particles locatedat austenite LABs. Cold rolled duplex stainless steelUranus 45N was shown by neutron diffraction to haveferrite and austenite phase textures identical to those oftheir single phase steel counterparts. In fact, a loss of theOR between ferrite and austenite was found with EBSD,due to complex deformation structures imaged in theTEM (Baudin et al., 2002). Electron backscatter diffr-action may also suggest ORs that are useful for global
texture calculations to be coupled to hot stage XRDtexture characterisation of the product phase [e.g. ina low carbon steel (Bruckner and Gottstein, 2001)]. In(Ni,Co)O single crystals, neutron diffraction gaveaverage results on crystal quality (here, mosaicity),XRD gave details on lattice parameters and EBSDprovided spatial information on crystal orientation(Brewer et al., 2002). The ORs between Ni and NiOwere obtained by carrying out EBSD (sensitive to NiOonly) and XRD (sensitive to the Ni substrate) in thesame area (Woodcock et al., 2004).
As XRD and TEM are well known techniques theycan easily be combined to get quantitative data oncrystal structure and orientation. Neutron diffraction ismuch less commonly used but provides useful bulkcharacterisation. Transmission electron microscopy isstill necessary for fine scale imaging and crystal charac-terisation, although sample preparation is destructiveand more tedious and analysed areas are small.
In situ investigations
Very few EBSD analyses associated with in situ inves-tigations have been published yet. Hot stage microscopyassociated with crystallographic EBSD characterisationof initial and/or final state has already provided newinteresting data (Table 28). With the development ofhigh speed, highly sensitive EBSD systems this is a
a
b
a TEM thin foil observation of heavily dislocated,slightly misoriented bainite ‘b’ next to a ferrite grain ‘f’together with retained austenite particles ‘a’; b EBSDmap of a ferrite grain: angular deviation from the crys-tal orientation of the grain centre (from 0 in white up to4u in black), showing bainite areas in dark
16 Coupled use of EBSD and TEM to characterise bai-
nite in a low alloy TRIP aided steel: hatched particles
are retained austenite, other ferrite grains are in
white. After Zaefferer et al. (2004)
a
b
a information obtained with TEM reported on a micro-graph; b EBSD map of misorientation with respect topoint X
17 Coupled EBSD and TEM investigation of the internal
structure of plate martensite in Fe–32.85 at.-%Ni: a
misorientation profile is superimposed, showing mis-
orientation up to 3u, following the change in OR
between austenite (A) and martensite (M). After
Shibata et al. (2005)
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
104 International Materials Reviews 2007 VOL 52 NO 2
promising way to get information about phase trans-formation mechanisms, at least at the sample surface.Environmental SEMs may be equipped with EBSDanalysis. Nothing seems to have been published yetabout in situ EBSD investigation of surface reactions butthe feasibility of EBSD analysis in various environ-mental and thermal conditions has already beendemonstrated in water and nitrogen environments forboth conducting and non-conducting materials(Garmestani and Harris, 1999; Habesch, 2000). In situphase transformation of low alloy TRIP aided steelsconfirmed that retained austenite gradually transformedupon straining, coarse equiaxed particles being lessstable that elongated ones in the area monitored byEBSD (Oh et al., 2002; Park et al., 2002).
Electron backscatter diffraction in hot stage SEM hasbeen made possible by a low effect of temperature onpattern quality for various materials up to 650 or even700uC (Garmestani and Harris, 1999; Seward et al.,2002) and by the recent development of fast EBSDsystems (Wright et al., 2005). Published studies mainlyconcentrated on recrystallisation and grain growth, e.g.in aluminium (Ferry, 1998; Huang et al., 2002) andnickel and deformed NaCl (Piazolo et al., 2004a). Somestudies have already been published on phase transfor-mations, e.g. in hcp a into bcc b and then back into hcpa transformations of titanium (Seward et al., 2002). Byheating titanium, bcc b phase transformed from hcp amay appear as plates with a Burgers OR and a tentshaped surface relief (Fig. 18) (Seward et al., 2004).Orientation relationships close to NW were identified inhcp a into fcc b and then back into hcp a transforma-tions of cobalt (Wright et al., 2005). Phase transforma-tions in iron oxides such as that from hematite intomagnetite have also been investigated using hot stageEBSD (Piazolo et al., 2004b). The main difficulties inhot stage EBSD mapping are temperature measurements(Seward et al., 2002) and the high kinetics of phasedevelopment, which may restrict in situ monitoring toonly a small area, possibly leading to sampling effects ifno repeated experiments are to be made.
Three-dimensional considerations
Electron diffraction cannot give information aboutcrystal orientation in the bulk of the material; 3D-XRD techniques are being developed for this pur-pose but still require a high flux synchrotron X-rayfacility to get the 3D shape of individual crystals andof orientation gradients (Fu et al., 2003; Barabashet al., 2006). The spatial resolution is coarser (abouta few mm) than that of EBSD but the angular re-solution is high (y0.05u) (Juul Jensen, 2000). Beingnon-destructive, 3D-XRD may be used for in situ inves-tigation of e.g. recrystallisation (Juul Jensen, 2000)and possibly phase transformations. However, it is farfrom being readily available, so that other methodsare generally used to get 3D information about crystalmorphology.
Three-dimensional investigations
The reconstruction of the 3D morphology is difficult(Fig. 19) and involves both careful experimental section-ing and complex 3D image processing. For coarsegrained materials, it may be useful to analyse bothparallel sides of the sample and to assume that no othercrystal is embedded in the bulk [e.g. in single shear lapT
ab
le28
Investi
gati
on
of
ph
ase
tran
sfo
rmati
on
sb
yE
BS
Dasso
cia
ted
wit
hin
sit
um
icro
sco
py
Mate
rial
Ho
tsta
ge
devic
eIn
form
ati
on
pro
vid
ed
by
ho
tsta
ge
devic
eIn
form
ati
on
pro
vid
ed
by
EB
SD
Ref.
Sp
lat
coolin
gand
melt
sp
innin
gof
molten
Si
Infr
are
dm
icro
scop
yS
pre
ad
ing
of
conta
ct
zone
The
conta
ct
zone
consis
tsof
fine,
eq
uia
xed
gra
ins
Nag
ashio
et
al.,
2004;
Nag
ashio
and
Kurib
ayashi,
2006
Form
ation
of
prim
ary
and
Wid
mansta
tten
ferr
ite
inlo
wcarb
on
ste
el
Laser
confo
cal
mic
roscop
yTim
ere
solv
ed
monitoring
nucle
ation
and
gro
wth
of
ind
ivid
ualp
hases
Mis
orienta
tion
betw
een
prim
ary
and
Wid
mansta
tten
ferr
ite
Phela
nand
Dip
penaar,
2004;
Phela
net
al.,
2005
Fe–32
at.
-%N
ib
icry
sta
lsLig
ht
op
ticalm
icro
scop
yM
onitoring
revers
etr
ansfo
rmation
from
mart
ensite
into
auste
nite
Fin
alm
icro
textu
re:
back
toorig
inalauste
nite
cry
sta
lorienta
tions
tog
eth
er
with
form
ation
of
LA
Bs
Ued
aet
al.,
2004
Fe–4. 1
8at.
-%C
rLig
ht
op
ticalm
icro
scop
yN
ucle
ation
sites
of
auste
nite
during
heating
Eff
ect
of
pare
nt
GB
mis
orienta
tion,
sub
seq
uent
ab
norm
al
gro
wth
and
twin
nin
gof
auste
nite
Wata
nab
eet
al.,
2004
and
2005b
Cu–A
l–B
eS
MA
X-r
ay
diffr
action*
Inte
rnalstr
esses
(one
sele
cte
dg
rain
)C
rysta
lorienta
tion
of
neig
hb
ouring
gra
ins
Kaouache
et
al.,
2004
*Tensile
sta
ge
devic
e.
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 105
Sn–3.5Ag solder joints (Telang and Bieler, 2005a)].Habit planes of plates may be determined by analysis oftwo perpendicular sections (Fig. 19c) such as in the bto a’’ phase transformation of titanium based b-Cezalloy (Zimmermann and Humbert, 2002). Electronbackscatter diffraction analysis of interfaces and facets
has been recently reviewed (Randle, 2004) and will notbe detailed here.
Another way to get 3D information on phasemorphology and connectivity is deep etching performedafter EBSD analysis (Fig. 19b). The morphology ofeutectic silicon was studied as a function of alloyingelements and local crystallography of phases in Al–Sihypoeutectic alloys (Nogita and Dahle, 2001a and2001b). Proeutectoid cementite in Fe–1.34C–13.1Mnwas shown to appear as both monolithic single crystalsand polycrystalline conglomerates of parallel laths, theOR with the austenite matrix being different from onecase to the other (Mangan et al., 1999). The morpho-logy of intergranular cementite in Fe–1.3C–13Mn wasstudied by TEM observation of deep etched materialand by EBSD investigation of OR with the austenitematrix (Kral and Spanos, 2003). Deep etching is alsouseful to investigate 3D connectivity of phases, e.g. ineutectic M7C3 carbides of white cast irons. These appearas interconnected colonies of parallel rods. Colonies aresingle crystalline in hypereutectic alloys and polycrystal-line in hypoeutectic alloys (Randle and Powell, 1993).3D analysis also gives access to nucleation sites, forexample in successive nucleation of MnS particles atoxides, of VC carbides at MnS sulphides, and of pearliteat VC carbides in intragranular pearlite of hypereutectic0.8C–12Mn–V steel (Guo et al., 2002).
The surface geometry may be determined by use ofstereo pairs as is already done on fracture surfaces(Hebesberger et al., 2000). Provided that the surface
(a)
(b)
(c)
a backscattered electron image (A1, A2, A3: b allotrio-morphs; P1: intragranular b); b pole figures of {111}band {11.0}a; c pole figures of {110}
band {00.1}a for
grains A3 and a218 In situ EBSD investigation of ORs between growing
allotriomophic b phase and starting a phase at 882uC:
after Seward et al. (2004)
(a)
(c)
(b)
(d)
(e)
19 3D morphological investigations to assess the hidden
volume a of the iceberg by b deep etching, c non-
parallel plane sectioning, d serial sectioning and e
one plane trace analysis: a practical example is given
in Fig. 6
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
106 International Materials Reviews 2007 VOL 52 NO 2
relief does not impede EBSD analysis by shadowingeffects, phase identification on fracture surface ofmultiphase materials [e.g. in c za2 TiAl alloys(Hebesberger et al., 2000)] is possible. Grain boundarygeometry (Randle, 2004) provides additional data forthe understanding of segregation phenomena [e.g. inNiz7 at. ppm S (Cornen and Le Gall, 2004)] andprecipitation.
The most widely used technique to get 3D informa-tion from EBSD is serial sectioning (Fig. 19d) either toget the 3D morphology alone (‘uncoupled’ use inTable 29) or even to get an EBSD map of all or partof the slices (‘coupled’ use in Table 29). Manual orautomatic metallographic polishing is the most conven-tional method. Ion milling is now in development (Chenand Rudolph, 2003). Dual beam SEMs now allow moreor less automated sequences of FIB milling followed byEBSD analysis in the same apparatus (Groeber et al.,2006; Zaafarani et al., 2006; Konrad et al., 2006). As thespecimen has to be moved between successive millingand EBSD mapping sequences, absolute spatial locationand orientation of the sample must be adjusted for eachslice (King et al., 2000) and distortion due to, e.g.electron beam drift during EBSD mapping is to bestrictly avoided. Data processing is now much easier,thanks to the increase in computer power and thedevelopment of dedicated software.
Three-dimensional analysis from 2D results
Just as for exposed versus non-exposed facets (Randle,2004), many EBSD investigations rely on analysis oftraces (Fig. 19e) to determine facet orientation in crystalgrowth, interphase interfaces, habit planes and interac-tions between cracks and microstructure of productphases (Table 30) for both metallic and non-metallicmaterials. Most studies determine probable crystal planefamilies to which the observed traces could belong.Several methods exist to assess the reliability of theseresults. If crystal planes are considered, one cannormalise the results (e.g. fraction of interfaces whosetraces are closer to experimental ones than a user definedthreshold) by the frequency that would be obtained ifthese planes were randomly oriented with respect to thecrystal. This is particularly useful if the plane multi-plicity is high due to crystal symmetry. For example, in asimulated HAZ of a low carbon steel, {223} and {557}austenite planes are of high multiplicity (12) but ‘habit’traces of bainite lath groups were still two times morefrequently close to the trace of these planes than arandomly oriented plane would have been (Gourgues,2003). The same kind of calculations was carried out forthe growth direction of bainite in a 0.6C–1.5Si–1Mnsteel (Cabus, 2005). Such normalising takes the actualcrystallographic orientation of phases into account andthus reduces the bias induced by the crystallographictexture. Habit planes are generally expressed in theframe of the parent phase, so that normalising has to becarried out for each grain orientation of the parentphase separately, which can be tedious when this phaseis totally absent from the resulting microstructure.
Stereological methods based on extensive data analy-sis over a number of non-parallel sections are nowavailable to up to date computer systems and providestatistically based information on interfaces [e.g. in GBsegregation of Nb in TiO2 (Pang and Wynblatt, 2006)].Other methods assuming tetrakaidecahedron shaped
grains help getting from 2D sections to 3D investiga-tions of ‘effective’ grain size after phase transformationin hot rolled steel (Bhattacharjee and Davis, 2002).Another way to investigate the 3D connectivity of lowmisoriented phases is to break the sample by brittle‘crystallographic’ fracture (e.g. by cleavage in bcc steels)and to look at fracture surfaces. Although lower bainiteand martensite in steel HAZs come with a high den-sity of HABs in 2D sections, areas of close crystalorientations within a given parent grain are actuallyconnected to each other in 3D, so that the size ofcleavage facets is much coarser (in this case, close to thehuge austenite grain size) and the toughness is muchlower than predicted from the ‘crystallographic’ grainsize calculated from individual 2D EBSD sections,although embedded misoriented crystals induce someroughness in cleavage fracture surfaces (Gourgues,2003).
In summary, although the vast majority of EBSDresults focus on 2D considerations only, one must notforget how to get to a 3D view, which is relevant tomaterial properties and reduces sampling effects such aspseudo variant selection. This is particularly true forcomplex shaped microstructures such as bainite andmartensite, and for all studies involving nucleation orvariant selection. The fast development of new 3Danalysis techniques should make it possible to increas-ingly include 3D considerations into the crystallographyof many phase transformations.
Coupling EBSD results with quantitativemodellingMany modelling approaches now use EBSD data as anexperimental basis. Some examples are given in thissection to illustrate the specific use of EBSD for thatpurpose.
Modelling phase transformations
Geometry considerations have been used for a decade tointerpret EBSD results of interphase interface proper-ties, phase connectivity and stereology (Table 31). Suchmodels should now benefit from 3D investigations asmentioned in the previous section. Another field ofmodelling is coupling EBSD data with heat and/or fluidflow calculations, especially for solidification (Table 32).In many instances, heat and fluid flow models aresimplified either to speed up calculations or to avoidintroduction of parameters that are difficult to calibrate(e.g. thermal exchange coefficients).
Little literature is available concerning couplingEBSD data and modelling of phase transformationkinetics, except for CET (see section on solidification).A simple model for diffusion anisotropy was usedto discuss results of oxygen tracer diffusion inLa2–xSrxCuO4¡d (Claus et al., 1996). Grain boundarymigration coupled with diffusion anisotropy allowedto explain EBSD data on heterogeneous kinetics ofdiscontinuous ordering in Fe–50 at.-%Co (Bischoffet al., 1998). Simple analytic kinetic models were usedand their results compared with crystal connectivitydetermined with EBSD to investigate fragmentation ofthe nickel skeleton in containerless eutectic solidificationof Ni–18.7 at.-%Sn (Li et al., 2005b), oscillation ofperitectic solidification in Sn–1.4 wt-%Cd (Zeisler-Mashl and Lograsso, 1997) and dendrite tip under-cooling in laser treated Fe–Cr–Ni alloys (Fukumoto and
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 107
Tab
le29
3D
mic
rostr
uctu
ral
investi
gati
on
su
sin
gE
BS
D
Mate
rial
Ph
ase
tran
sfo
rmati
on
Co
up
led
or
un
co
up
led
EB
SD
an
aly
sis
Resu
lts
(new
info
rmati
on
)R
ef.
Al–
10S
iE
ute
ctic
solid
ific
ation
Uncoup
led
Rad
ialg
row
thof
eute
ctic
Sifr
om
Sip
oly
hed
raD
ahle
et
al.,
2005
Al–
1S
i–0. 1
7Ti
Solid
ific
ation
(feath
ery
gra
ins)
Coup
led
Dend
rite
gro
wth
direction,
twin
nin
gp
lanes
Henry
et
al.,
1998
2618
alu
min
ium
allo
yS
olid
ific
ation
(thix
ofo
rmin
g)
Uncoup
led
Glo
bula
rm
orp
holo
gy
as
alread
ysug
geste
db
y2D
analy
sis
,H
AB
sX
iaand
Tausig
,1998
AL–6X
Nsup
era
uste
nitic
sta
inle
ss
ste
el
Inte
rgra
nula
rp
recip
itation
of
sp
hase
Coup
led
(EB
SD
every
tenth
slic
e)
Morp
holo
gy,
cry
sta
llog
rap
hy,
GB
pla
ne,
OR
betw
een
sand
matr
ixLew
iset
al.,
2006
Auste
nitic
sta
inle
ss
ste
el
for
hyd
rog
en
refo
rmin
gP
recip
itation
of
M23C
6C
oup
led
All
cre
ep
cavitie
sare
connecte
dto
M23C
6;
these
M23C
6
are
less
freq
uently
incub
e–cub
eO
Rw
ith
the
matr
ixth
an
ifnot
connecte
dto
cre
ep
cavitie
s
Ab
dulW
ahab
and
Kra
l,2005
Cu–4
wt-
%Ti
Dis
continuous
pre
cip
itation
from
ain
toaz
Cu
4Ti
Coup
led
Exte
nsiv
ein
form
ation
ab
out
morp
holo
gy
and
cry
sta
llog
rap
hy
of
colo
nie
sM
ang
an
and
Shifle
t,1997
Fe–C
pla
incarb
on
ste
els
Eute
cto
idp
earlite
form
ation
Coup
led
3D
cry
sta
llog
rap
hy
of
pearlite
colo
nie
s,
OR
betw
een
ferr
ite
and
cem
entite
linked
tonucle
ation
sites
Mang
an
and
Shifle
t,1999
Fe–0. 8
C–12M
n–0. 3
VE
ute
cto
idp
earlite
form
ation
Uncoup
led
3D
imag
ing
:nucle
ation
cond
itio
ns;
EB
SD
:no
OR
betw
een
auste
nite
and
pearlitic
ferr
ite
Guo
et
al.,
2002
Fe–0. 8
C–12. 3
Mn
Form
ation
of
Wid
mansta
tten
cem
entite
pla
tes
Coup
led
Inte
rlocked
pla
tes
are
infa
ct
diffe
rent
cry
sta
lsand
not
bifurc
ations
of
the
sam
ecry
sta
l;in
form
ation
ab
out
nucle
ation,
gro
wth
and
OR
s
Mang
an
et
al.,
1997
Low
carb
on
ste
el
Tra
nsfo
rmation
of
auste
nite
into
inte
rgra
nula
rfe
rrite
Coup
led
Poly
cry
sta
lline
ferr
ite
(inclu
din
gLA
Bs);
all
cry
sta
lsare
inconta
ct
with
aM
nS
nucle
ating
part
icle
;th
enum
ber
of
ferr
ite
cry
sta
lsp
er
MnS
incre
ases
with
the
siz
eof
the
MnS
part
icle
Yokom
izo
et
al.,
2003
Hig
hp
urity
Fe–0. 1
2C
–3. 2
8N
iTra
nsfo
rmation
of
auste
nite
into
inte
rgra
nula
rp
rim
ary
and
Wid
mansta
tten
ferr
ite
Coup
led
Peak
morp
holo
gy:
pyra
mid
alsin
gle
cry
sta
ls;
lath
sand
second
ary
pla
tes:
poly
cry
sta
lsw
ith
inte
rnalLA
Bs
(rep
eate
dnucle
ation);
one
cry
sta
lorienta
tion
far
from
auste
nite
GB
s
Sp
anos
et
al.,
2005;
Kra
land
Sp
anos,
2005
Hig
h-p
urity
Fe–0. 1
2C
–3. 2
8N
iTra
nsfo
rmation
of
auste
nite
into
ferr
ite
sid
ep
late
sU
ncoup
led
Prim
ary
ferr
ite
film
and
sid
ep
late
ssep
ara
ted
by
aLA
B:
there
isin
deed
nucle
ation
of
Wid
mansta
tten
ferr
ite,
not
only
unsta
ble
gro
wth
of
the
ferr
ite
film
Sp
anos
and
Hall,
1996
Auste
mp
ere
dd
uctile
cast
iron
Tra
nsfo
rmation
of
auste
nite
into
ausfe
rrite
Coup
led
Fatig
ue
cra
cks
arr
est
at
packet
bound
aries;
sta
ge
Icra
cks
are
conta
ined
in{1
11}
pla
nes
of
auste
nite
Marr
ow
et
al.,
2002
0. 0
7C
–0. 8
Mn–3. 5
Ni–
1. 6
Cu–0. 6
Mo
(HS
LA
–100)
Tra
nsfo
rmation
of
auste
nite
into
‘coars
e’m
art
ensite
Coup
led
Tap
ere
d3D
morp
holo
gy
and
cry
sta
lin
dic
es
of
inte
rface
pla
nes
Row
enhors
tet
al.,
2006
Ni–
45. 6
at.
-%A
lIn
terg
ranula
rseg
reg
ation
Uncoup
led
The
GB
energ
yd
ep
end
son
the
GB
pla
ne
even
for
rand
om
GB
sA
mouyalet
al.,
2005
Dia
mond
on
sili
con
Cry
sta
lg
row
thC
oup
led
,p
ara
llel
toth
efilm
pla
ne
Colu
mnar
{111}
gro
wth
Chen
and
Rud
olp
h,
2003
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
108 International Materials Reviews 2007 VOL 52 NO 2
Tab
le30
Tra
ce
an
aly
sis
fro
mE
BS
Dre
su
lts
(excep
tth
ose
co
ncern
ing
the
gro
wth
dir
ecti
on
inso
lid
ificati
on
,see
co
rresp
on
din
gsecti
on
)
Mate
rial
Ph
ase
tran
sfo
rmati
on
Facet
An
aly
sis
co
nd
itio
ns
Resu
lts
Ref.
Pb
(Mg
0. 3
3N
b0. 6
7)O
3z
Pb
TiO
3R
eactive
sin
tering
Cry
sta
lfr
ee
surf
aces
Tra
ce
analy
sis
(no
deta
ilp
rovid
ed
){1
00}
facets
Liet
al.,
1998
Kaolin
ite
and
dic
kite
min
era
lsC
rysta
lg
row
thC
rysta
lfr
ee
surf
aces
Tra
ce
analy
sis
Orienta
tion
of
late
ralfa
cets
Kog
ure
et
al.,
2005;
Kam
ed
aet
al.,
2005
SrT
iO3
inTiO
2rich
liquid
Coars
enin
gof
cry
sta
lsC
rysta
lfr
ee
surf
aces
Tra
ce
analy
sis
zste
reolo
gic
alm
od
el
Shap
echang
ein
to{1
00}
facets
and
,100
.zone
axes
during
coars
enin
gS
ano
et
al.,
2005
2. 2
5C
r–1M
oste
el
Auste
nite
into
bain
ite
Cle
avag
ecra
cks
Tra
ces
plo
tted
onto
ste
reog
rap
hic
pro
jection
Com
patib
lew
ith
{100}
pla
nes
of
bain
ite
Bouyne
et
al.,
1998
Thre
elo
wcarb
on
ste
els
Auste
nite
into
acic
ula
rfe
rrite
and
bain
ite
Cle
avag
ecra
cks
Tra
ce
analy
sis
,cra
ck
arr
est
cond
itio
ns
{100}
traces,
cra
ck
devia
tion
for
mis
orienta
tion
ang
le.
15u
Dıa
z-F
uente
set
al.,
2003
Dup
lex
sta
inle
ss
ste
el
Ferr
ite
into
auste
nite
Fatig
ue
cra
cks
Tra
ce
analy
sis
Str
ong
eff
ect
of
the
two
phase
mic
rostr
uctu
reon
cry
sta
lorienta
tion
of
traces
Gourg
ues
et
al.,
2004
0. 7
8C
mic
roallo
yed
ste
el
Auste
nite
into
pearlite
Cle
avag
ecra
cks
Orienta
tion
units
(5sin
gle
cry
sta
lfe
rrite
inp
earlite
colo
nie
s)
The
siz
eof
orienta
tion
units
isth
esam
eas
that
of
cle
avag
efa
cets
corr
ecte
dfr
om
2D
/3D
ste
reolo
gy
Cotr
ina
et
al.,
2003
Auste
mp
ere
dd
uctile
cast
iron
Auste
nite
into
ausfe
rrite
Fatig
ue
cra
cks
2D
and
3D
(serialsectionin
g)
pro
pag
ation
pla
nes
{111}
pla
nes
of
pare
nt
(and
reta
ined
)auste
nite
Marr
ow
et
al.,
2002
Mo
5S
i 3(C
r,Ti)
Solid
ific
ation
(seg
reg
ate
dzones)
Cle
avag
enear
ind
enta
tion
mark
sTra
ce
analy
sis
(001)
cry
sta
lp
lanes
Str
om
and
Zhang
,2005
Hig
hstr
eng
thlo
wallo
yste
el
Auste
nite
into
up
per
bain
ite
Str
aig
ht
inte
rfaces
of
lath
gro
up
sTra
ce
analy
sis
{557}
and
{223}
auste
nite
pla
nes
more
pro
bab
leth
an
{111}
auste
nite
pla
nes
Lam
bert
-Perlad
eet
al.,
2004a
0. 8
C–12. 3
Mn
ste
el
Auste
nite
into
pearlite
Hab
itp
lanes
Work
on
EB
SD
patt
ern
s(0
01) c
//{1
25} c
with
Pitsch–P
etc
hO
R,
(101) c
//{2
11} c
with
Isaic
hev
OR
;(0
01) c
//{1
21} c
with
Bag
ary
ats
ky
OR
dep
end
ing
on
the
ind
ivid
ualcolo
ny
Mang
an
and
Shifle
t,1998
Cu–3
wt-
%Ti
Cellu
lar
pre
cip
itation
Hab
itp
lanes
Work
on
EB
SD
patt
ern
s(1
11) a
//(0
10) b
Mang
an
and
Shifle
t,1998
Ti
hcp
ain
tob
cc
bH
ab
itp
lane
of
bIn
situ
trace
analy
sis
zshap
econsid
era
tions
Near
to{3
34} b
,conta
ins
the
[0001] a
//[1
10] b
Sew
ard
et
al.,
2004
Hig
hstr
eng
thlo
wallo
yste
el
Auste
nite
into
bain
ite
Str
aig
ht
inte
rface
of
lath
gro
up
sTra
ce
analy
sis
{557}
or
{223}
but
not
{111}
pla
nes
of
auste
nite
Gourg
ues,
2003
Ti–
25A
l–24
at.
-%N
ba
2in
toO
GB
pla
nes
Auto
mate
dtr
ace
analy
sis
of
MR
and
inte
rphase
inte
rfaces
(110)
for
39%
of
twin
rela
ted
cry
sta
lsLiet
al.,
2004
TiO
2d
op
ed
with
Nb
GB
seg
reg
ation
GB
pla
ne
Tra
ce
analy
sis
zste
reolo
gic
alm
od
el
The
am
ount
of
seg
reg
ate
dN
bis
invers
ely
pro
port
ional
toth
efr
eq
uency
of
GB
s;
Nb
seem
sto
flatt
en
the
GB
energ
yd
istr
ibution
Pang
and
Wynb
latt
,2006
0. 6
C–1. 5
Si–
1. 5
Mn
ste
el
Auste
nite
into
bain
ite
Str
aig
ht
inte
rfaces
of
lath
gro
up
sTra
ce
analy
sis
znorm
alis
ing
66%
of
traces
at
,10u
from
,110
.p
roje
ction
onto
sam
ple
surf
ace
(50%
for
rand
om
lyoriente
dp
lanes)
Cab
us,
2005
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 109
Tab
le31
Mo
dellin
gp
hase
geo
metr
yu
sin
gE
BS
Dd
ata
Ph
ysic
al
ph
en
om
en
on
of
inte
rest
Mate
rial
Mo
del
Use
of
or
co
mp
ari
so
nw
ith
EB
SD
data
Ref.
Hete
roep
itaxy
ZrO
2/Y
2O
3onto
Al 2
O3
Coin
cid
ence
site
latt
ice
bound
aries
More
OR
sp
red
icte
dth
an
those
ob
serv
ed
by
EB
SD
Cain
and
Lang
e,
1994;
Cain
et
al.,
1995
Directionalsolid
ific
ation
CM
SX
–4
and
CM
186LC
Ni
base
sup
era
lloys
Coin
cid
ence
site
latt
ice
bound
aries
With
the
ob
serv
ed
,001
.fib
rete
xtu
re,
not
all
pre
dic
ted
coin
cid
ence
site
latt
ice
bound
aries
are
ob
serv
ed
Ard
akaniet
al.,
2000
Fre
esurf
ace
energ
yS
rTiO
3In
terp
reta
tion
of
results
of
ato
mic
forc
em
icro
scop
yR
esults
were
coup
led
toE
BS
Dcry
sta
lorienta
tion
data
tod
ete
rmin
eanis
otr
op
icfr
ee
surf
ace
energ
yS
ano
et
al.,
2003a
Solid
sta
tesp
read
ing
Cop
per
onto
cop
per
Em
bed
ded
ato
mm
od
el
For
tim
es
up
to30
ps,
the
mod
elw
ork
sw
ell
for
morp
holo
gy
of
phases,
but
not
with
cry
sta
lalig
nm
ent
ob
serv
ed
by
EB
SD
Read
et
al.,
2004
Rela
tive
HA
Benerg
yN
irich
NiA
lE
mb
ed
ded
ato
mm
od
el
EB
SD
zscannin
gp
rob
em
icro
scop
yg
ive
valu
es
of
0. 2
up
to1. 1
;th
em
od
elg
ives
0. 3
up
to0. 9
Am
ouyalet
al.,
2005
GB
netw
ork
No
part
icula
rm
ate
rial
Geom
etr
yof
GB
netw
ork
How
toextr
act
the
rand
om
GB
netw
ork
from
EB
SD
data
Kin
get
al.,
2000
GB
wett
ing
and
finite
siz
escalin
gG
ain
Zn
Latt
ice
mod
elfo
rsam
plin
geff
ects
sim
ilar
toth
ose
of
EB
SD
measure
ments
Eff
ect
of
sam
ple
wid
thon
the
dep
thof
inte
rgra
nula
rin
vasio
nof
Zn
by
Ga
Tra
skin
eet
al.,
2005
Variant
sele
ction
Bain
ite
inste
els
Cro
ss-s
ection
eff
ects
on
textu
red
ete
rmin
ation
(i.e
.p
seud
ovariant
sele
ction)
Even
while
takin
gp
seud
ovariant
sele
ction
into
account,
avera
ge
textu
res
still
show
variant
sele
ction
Cab
us,
2005
Directionalsolid
ific
ation
Ni–
Cr–
Co–Ti–
Al–
Mo–S
isup
era
lloy
Eff
ect
of
textu
restr
eng
thenin
gon
gra
inclu
ste
ring
EB
SD
measure
dg
rain
clu
ste
ring
ag
rees
with
geom
etr
iceff
ects
ind
uced
by
textu
restr
eng
thenin
gW
est
and
Ad
am
s,
1997
Tab
le32
Co
up
lin
gE
BS
Dd
ata
wit
hh
eat/
flu
idfl
ow
calc
ula
tio
ns
Ph
ysic
al
ph
en
om
en
on
of
inte
rest
Mate
rials
an
dp
rocesses
Mo
del
Use
of
EB
SD
data
Ref.
Menis
cus
shap
eS
lab
casting
of
ultra
low
carb
on
ste
els
Bic
kerm
an
(analy
tical)
fluid
flow
eq
uations
The
actu
alshap
eof
frozen
menis
cus
dete
rmin
ed
with
EB
SD
ag
rees
with
mod
eloutp
ut
Seng
up
taet
al.,
2006
Direction
of
dend
rite
gro
wth
Tw
inro
llcasting
of
Fe–3S
i3D
dend
ritic
gro
wth
with
sim
plif
ied
geom
etr
y,
takin
gfluid
flow
into
account
EB
SD
dete
rmin
ed
density
of
nucle
iin
chill
zone
isa
mod
elin
put;
the
dend
rite
gro
wth
direction
(mod
eloutp
ut)
com
pare
sw
ell
with
that
measure
dw
ith
EB
SD
Takata
niet
al.,
2000
Form
ation
of
feath
ery
gra
ins
during
solid
ific
ation
Direct
chill
casting
of
AA
1050
alu
min
ium
allo
yin
gots
Flu
idflow
inm
ould
:eff
ect
of
fluid
convection
Localm
icro
textu
rep
red
icte
db
yth
em
od
elis
com
pare
dto
EB
SD
data
Henry
et
al.,
2004
Com
petitive
gra
ing
row
thd
uring
solid
ific
ation
Directionalsolid
ific
ation
of
IN738LC
,IN
718
and
sin
gle
cry
sta
lnic
kelb
ase
sup
era
lloys,
solid
ific
ation
of
Zn
(hot
dip
galv
anis
ing
)
Heat
transfe
r,d
end
rite
or
gra
ing
row
th(c
ellu
lar
auto
mato
nz
finite
ele
ment
analy
sis
),coup
led
or
not
with
tem
pera
ture
calc
ula
tions
Pre
dic
ted
localte
xtu
re,
gra
insiz
eand
shap
eare
com
pare
dw
ith
EB
SD
data
for
pro
cess
op
tim
isation
Kerm
anp
ur
et
al.,
2000;
Cart
er
et
al.,
2000;
Sem
oro
zet
al.,
2002b
;N
ew
ell
et
al.,
2005;
Xu
et
al.,
2002a
and
2002b
Colu
mnar
toeq
uia
xed
transitio
nand
str
ay
gra
info
rmation
insolid
ific
ation
Weld
ing
and
laser
meta
lfo
rmin
gof
nic
kelb
ased
sup
era
lloys;
hot
dip
galv
anis
ation
with
Zn
allo
y
Sim
plif
ied
therm
alm
od
el(e
.g.
Rosenth
al
eq
uations
for
weld
ing
)g
ives
coolin
gra
teand
tem
pera
ture
gra
die
nt;
akin
etic
mod
elis
used
inp
ost-
pro
cessin
gcalc
ula
tions
Density
of
nucle
i(inp
ut
for
the
mod
el);
the
pre
dic
ted
volu
me
fraction
of
eq
uia
xed
gra
ins
iscom
pare
dto
EB
SD
ob
serv
ations
Gaum
ann
et
al.,
2001;
Vitek
et
al.,
2004;
Quirog
aet
al.,
2004
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
110 International Materials Reviews 2007 VOL 52 NO 2
Kurz, 1997). Agglomeration of silicon particles in semi-solid Al–Si alloy evidenced by EBSD investigations wasalso suggested by the failure of a liquid film migrationmodel to fit experimental data (Hogg and Atkinson,2005).
Modelling local crystallography and variant selection
Crystallographic models of ORs have been only scarcelyused together with EBSD data, except for counting thenumber of possible variants or calculating the MRsbetween variants (for instance, the present author foundat least 11 EBSD related papers with detailed descrip-tion of MRs between variants for KS or NW ORs).
Phenomenological theories of martensite crystallogra-phy are widely used to describe crystallographic featuressuch as ORs, habit planes and macroscopic transfor-mation eigenstrains. The traditional way to comparepredictions with experiment is TEM investigation of thinfoils, but EBSD may provide statistically significantdatasets for both ORs and (if coupled with surfaceanalysis techniques such as near field microscopy)transformation eigenstrains. EBSD is much less suitedto habit plane determination unless ‘habit’ planes aredefined by straight interphase interfaces (and may not be‘true’ habit planes). Predictions of PTMCs have beencompared with EBSD measurements for habit planes ofa’’ phase in titanium based b-Cez alloy (in fact EBSDresults suggest that the highest value of lattice invariantshear should be taken) (Zimmermann and Humbert,2002). ‘Habit planes’ close to {223} or {557} austeniteplanes were found for groups of upper bainite laths in asimulated HAZ of low carbon steel (Lambert-Perladeet al., 2004a). Both a Pitsch type OR and a macro-scopic shear eigenstrain of 0.20 were determined in Cu–39.3 at.-%Zn (from bcc b to fcc a bainite, treated as 9Rmartensite in the PTMC) (Marukawa et al., 2000), aswere habit plane and transformation eigenstrain deter-mined in Ti–9 wt-%Mo alloy (from bcc b into orthor-hombic a’’ phase) (Guo et al., 2000). The OR versuslattice invariant shear mechanism was discussed inlenticular a’ martensite of Fe–32.85 wt-%Ni (Shibataet al., 2005).
Coupled use of EBSD results and modelling generallyaims at accounting for variant selection in solid statephase transformations. Various criteria have beendeveloped to model how a limited number of variantsappear under certain conditions due to local stress orstrain fields.
In the active slip model, common plane and directionbetween parent and product phases should describe theslip system (of perfect or even partial dislocations) ofhighest Schmid factor. This local Schmid factor isgenerally calculated using the applied loading witheither a Sachs assumption (homogeneous stress in thematerial) or a more or less constrained Taylor assump-tion (homogeneous strain in the material) (Table 33).Active slip models give general trends and often showlimited agreement with experiments, as local stress andstrain fields in polycrystalline materials strongly varyfrom one grain to another, especially at grain bound-aries where nucleation often occurs.
Another criterion is the total eigenstrain calculated forthe simultaneous formation of at least two variants. Thiscriterion is used to describe self-accommodation by localvariant selection within the parent grain. Very goodagreement was found between groups of variants T
ab
le33
Pre
dic
tio
ns
of
acti
ve
sli
pm
od
els
for
vari
an
tsele
cti
on
co
mp
are
dw
ith
EB
SD
exp
eri
men
tal
resu
lts
Mate
rial
Tra
nsfo
rmati
on
co
nd
itio
ns
OR
Slip
syste
msz
localis
ati
on
assu
mp
tio
nC
rite
rio
nR
esu
lts
Ref.
Ti–
6A
l–4V
bcc
bin
tohcp
aaft
er
com
pre
ssio
nB
urg
ers
{110},
111
.b
and
{112},
111
.bz
Sachs
Maxim
um
resolv
ed
shear
str
ess
Qualit
ative
ag
reem
ent
with
exp
erim
enta
lavera
ge
textu
reM
ousta
hfid
et
al.,
1997b
Cu–A
l–B
eallo
yb
1auste
nite
into
mart
ensite
(tensio
n)
24
variants
(OR
not
sp
ecifie
d)
Localstr
esses
from
ela
stic
str
ain
sm
easure
db
yX
RD
;slip
syste
ms
not
sp
ecifie
dS
chm
idfa
cto
rS
mall
valu
es
of
crite
rion
lead
tono
phase
transfo
rmation
Kaouache
et
al.,
2004
Fe–32
at.
-%N
ib
icry
sta
lsfc
causte
nite
into
lenticula
rb
cc
mart
ensite
(tensio
n)
NW
(24
coup
les
of
variantz
hab
itp
lane)
Macro
scop
icshear
direction
and
hab
itp
lane
norm
al
Min
imum
constr
ain
ed
macro
scop
icstr
ain
at
GB
Ag
reem
ent
with
exp
erim
ent
on
identity
of
sele
cte
dvariants
Ued
aet
al.,
2001a
Ste
elw
ith
bain
itez
reta
ined
auste
nite
fcc
auste
nite
(c)
into
bcc
bain
itic
ferr
ite;
hot
rolli
ng
red
uction
of
0. 2
at
750uC
Near
KS
{111},
110
.c,
fully
constr
ain
ed
Taylo
rm
od
el
Schm
idfa
cto
rFeasib
ility
dem
onstr
ate
d,
pro
mis
ing
results
God
et
et
al.,
2001
0. 6
Cb
ain
itic
ste
el
fcc
auste
nite
(c)
into
bcc
bain
itic
ferr
ite;
channeld
iecom
pre
ssio
nz
ste
pq
uenchin
g
Uniq
ue,
clo
se
toK
Sand
NW
{111},
110
.c,
fully
constr
ain
ed
Taylo
rand
Sachs
mod
els
Com
mon
clo
se-p
acked
pla
nes
para
llelto
pla
nes
of
activate
dslip
syste
ms
Only
part
ialag
reem
ent
with
exp
erim
ent
Cab
us
et
al.,
2004a
Low
allo
yTR
IPaid
ed
bain
itic
ste
el
fcc
auste
nite
(c)
into
bcc
bain
itic
ferr
ite;
25–50%
red
uction
by
hot
rolli
ng
at
750uC
NW
,K
S{1
11},
110
.cz
corr
esp
ond
ing
Shockle
yp
art
ial
dis
locations;
Taylo
rm
od
el
Maxim
um
resolv
ed
shear
str
ess
‘Positiv
e’and
‘neg
ative’slip
variants
are
accounte
dfo
rb
ysp
read
ing
ab
out
cert
ain
NW
variants
Jonas
et
al.,
2005
Low
allo
yand
hig
hM
nTR
IPaid
ed
ste
els
fcc
auste
nite
(c)
into
bod
ycentr
ed
tetr
ag
onalor
hcp
mart
ensite
und
er
tensio
n
KS
and
Burg
ers
resp
ectively
{111},
110
.c,
Sachs
Assum
ption
Com
mon
clo
se
packed
pla
nes
para
llelto
pla
nes
of
activate
dslip
syste
ms
Good
ag
reem
ent
with
exp
erim
ent
on
identity
of
sele
cte
dvariants
God
et
et
al.,
2005
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 111
minimising the total eigenstrain (calculated with aPTMC) and groups actually found with EBSD in spearlike martensite of Cu–7.3Al–8.5 wt-%Mn (Wang et al.,2002), in plate martensite at a symmetrical 180u,211.
tilt boundary of a Fe–32 at.-%Ni bicrystal (Ueda et al.,2003a), in groups of three variants of hcp a phase intitanium (Wang et al., 2003). By contrast, lath marten-site subblocks evidenced with EBSD in Fe–C alloys werenot efficient self-accommodating structures according tothis criterion (Morito et al., 2003). Accommodation oflarge strains imposed by hot rolling was found to takeplace by variant selection in a low carbon bainitic steel(Matsuoka et al., 1999); however, in that study thelattice Bain strain was surprisingly chosen as thetransformation eigenstrain.
More complex criteria for strain accommodationwithin a given parent grain involve elastic interactionenergy calculated with the transformation eigenstrain(or more surprisingly lattice Bain strains in some cases)associated with candidate variants. Consistent resultsbetween observed variants and negative interactionenergy were found in austenitic AISI 301 austeniticstainless steel in tension and in compression, under theSachs assumption (Lee et al., 2005). By using a thres-hold value for the interaction energy criterion, goodaverage texture predictions were obtained for compactstrip processed high strength low alloy steel by usinganisotropic elasticity of phases and either the latticeBain strain (Humbert et al., 2002b) or the macroscopictransformation eigenstrain (Humbert et al., 2002a). InFe–32 at.-%Ni bicrystals, the interaction energy termcan be overwhelmed by the necessity for strain com-patibility at the GB according to the particular GBmisorientation (Ueda et al., 2001b and 2003a).
Coupled effects of transformation eigenstrain andphase geometry have been taken into account bymodelling local stress and strain fields based on theEshelby inclusion problem. The matrix is generallytaken to be the parent phase with isotropic oranisotropic elastic–plastic properties. Variants are mod-elled as oblate ellipsoids extending along their macro-scopic ‘habit’ plane. Mechanical coupling betweenphases is achieved under self-consistency assumptions.The model outputs are the interaction strain energy, andsometimes also the plastic strain in the parent phase as afunction of candidate variants. Such calculations werecompared to variant selection experimentally evidencedwith EBSD to account for the identity of variants instress free materials and as a function of applied stressfor, e.g. martensite formation in carburised steels(Karaman et al., 1998), for microtexture of upperbainite packets in low carbon steel HAZ (Gourgues,2003; Lambert-Perlade et al., 2004a) and for averagetexture affected by variant selection in Zircaloy–4 alloy(Humbert and Gey, 2003). There is still much to do inthis field, in particular to get accurate data on theconstitutive behaviour of phases under such extremeconditions (e.g. elevated temperature, possibly high localstrain rates and high strains). 3D EBSD coupled withlocal measurements of phase property (Zhao, 2006)could be a promising way to further improve thesemodels.
Modelling average transformation textures
The prediction of average transformation textures is ofutmost importance to infer product properties from
processing conditions. A lot of work has been carriedout in this field using EBSD data either alone or togetherwith XRD data (Table 34). Most studies in this fieldconcentrate either on variant selection effects on averagetexture or on accuracy of quantitative texture predic-tion. However, the texture of the parent phase isgenerally difficult to assess, even if some phase isretained, unless specific calculation methods are usedto retrieve it (see section on average textures in solidstate phase transformations).
Modelling resulting properties
In service properties of transformed microstructuresmay be tentatively predicted using EBSD data aboutmicrotexture, macrotexture and morphology of indivi-dual phases together with homogenisation models tocalculate ‘average’ or ‘macroscopic’ properties. Thesimplest homogenisation model used is the rule ofmixtures, which successfully predicted, e.g. the compres-sion peak stress of eutectic Al2O3–YAG–ZrO2 rods as afunction of the texture of individual phases (Murayamaet al., 2004b). More complex microtextures may inducefailure of this model to account for experimentalmeasurements, e.g. in wrought duplex stainless steels,where areas maintaining the KS OR between bothphases behave as ‘hard’ particles in a much softer matrix(Iza-Mendia et al., 1998).
More advanced micromechanical models have alsobeen used to predict yield properties of heterogeneousmaterials such as columnar and equiaxed zones of Ti–6Al–4V ingots under Taylor assumption and withcrystal plasticity models (Glavicic et al., 2003c and2004a). Local stresses arising from upset forging of Ti–6Al–4V were also calculated with a finite elementmethod, allowing local Taylor factors to be calculatedthanks to a simple micromechanical model (Bieler et al.,2005b) and cavity initiation sites to be related to localcrystal orientation and spheroidisation mechanisms(Bieler and Semiatin 2002, Bieler et al., 2005a and2005b). Micromechanical models of crystal plasticityhave been used to assess strain localisation in bainiticand martensitic microstructures of A508Cl.3 low alloysteel under tension (Sekfali et al., 2002) and yield andfracture properties of cast and aged duplex stainlesssteels (Bugat et al., 2001). In the two latter cases, localmeasurements of surface strains together with EBSDmapping of initial state were used to mesh the ‘actual’surface microstructure and to compare experimentalmeasurements of strains with model predictions. TheOR between parent and product phases, and MRbetween product variants play a key role in both sliplocalisation and cleavage cracking properties. Moresimply, the crystallographic packet size, which is theunit crack path for cleavage microcracking could also beused with a Hall–Petch type equation to predict thefracture toughness of low carbon acicular ferrite steels asa function of the packet size (Dıaz-Fuentes et al., 2003).It was also quantitatively used to predict fracturetoughness of a bainitic steel HAZ microstructure(Lambert-Perlade et al., 2004b). A local micromechani-cal model was used together with a transformationinduced superplasticity model to predict the evolution ofTi–6Al–4V/TiBw composite tensile loaded during ther-mal cycling around the matrix transformation tempera-tures (Schuh and Dunand, 2001).
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
112 International Materials Reviews 2007 VOL 52 NO 2
Tab
le34
Mo
dellin
gavera
ge
tran
sfo
rmati
on
textu
res
usin
gd
ata
gen
era
ted
wit
hE
BS
D
Mate
rial
Ph
ase
tran
sfo
rmati
on
Exp
eri
men
tal
data
Mo
del
Resu
lts
Ref.
Hig
hp
urity
iron
fcc
cin
tob
cc
aA
vera
ge
textu
re(b
oth
EB
SD
and
XR
D)
Sta
rtin
gfr
om
cte
xtu
re,
24
KS
variants
,no
variant
sele
ction
There
are
mis
sin
gvariants
inexp
erim
enta
lvers
us
pre
dic
ted
OD
F,
att
rib
ute
dto
gra
ing
row
thb
ut
not
tovariant
sele
ction
Ab
iko
et
al.,
2000
Low
carb
on
ste
el
bcc
ain
tofc
cc
Avera
ge
textu
reof
aand
caft
er
heat
treatm
ent
Sta
rtin
gfr
om
ate
xtu
re,
24
KS
variants
,no
variant
sele
ction
One
textu
recom
ponent
exp
erim
enta
llym
issin
g,
anoth
er
too
str
ong
vers
us
pre
dic
tions,
thus
there
isvariant
sele
ction
Park
et
al.,
2002
Low
carb
on
ste
el
fcc
cin
tob
cc
aA
vera
ge
textu
reS
tart
ing
from
typ
icalte
xtu
recom
ponents
of
defo
rmed
c,24
KS
variants
,no
variant
sele
ction
The
ap
hase
nucle
ate
sfr
om
defo
rmed
(not
recry
sta
llised
)c
Hurley
and
Hod
gson,
2001
Low
carb
on
ste
els
bcc
ain
tofc
cc
back
into
aE
BS
D:
gra
instr
uctu
re;
XR
D:
avera
ge
textu
reS
tart
ing
from
ate
xtu
re,
12
NW
variants
Pre
dic
tions
consis
tent
with
exp
erim
ent
for
one
ste
el,
insuff
icie
nt
textu
rem
em
ory
pre
dic
ted
for
two
oth
er
ste
els
Ryd
eet
al.,
1999
Low
carb
on
ste
el
bcc
ain
tofc
cc
EB
SD
:O
R;
XR
Dand
neutr
on
diffr
action:
avera
ge
textu
reS
tart
ing
from
cte
xtu
re,
24
KS
variants
;variant
sele
ction:
assum
ing
rolli
ng
str
ain
and
Taylo
rhyp
oth
esis
,b
oth
active
slip
and
transfo
rmation
work
crite
ria
were
teste
d
Mod
elp
red
ictions
consis
tent
with
exp
erim
ent
only
ifb
oth
crite
ria
are
com
bin
ed
Bru
ckner
and
Gott
ste
in,
2001
HS
LA*
ste
el(c
om
pact
str
ipp
rod
uction)
fcc
cin
tob
cc
aE
BS
D:
OR
;X
RD
:avera
ge
textu
reof
aand
reta
ined
cS
tart
ing
from
cte
xtu
re,
12
NW
variants
;variant
sele
ction:
ela
stic
str
ain
energ
y(B
ain
str
ain
or
macro
scop
ictr
ansfo
rmation
eig
enstr
ain
)
Possib
levariant
sele
ction:
pre
dic
ted
textu
res
are
sharp
er
than
exp
erim
enta
lones
Hum
bert
et
al.,
2002a
and
2002b
Low
allo
yTR
IPste
el
fcc
cin
tob
ain
itic
bcc
aE
BS
D:
OD
Fs
for
aand
reta
ined
c;X
RD
:volu
me
fraction
of
c
Thre
eB
ain
,12
NW
or
24
KS
variants
,no
variant
sele
ction
Pre
dic
ted
textu
rew
ith
rig
ht
com
ponents
but
too
weak
vers
us
exp
erim
ent
De
Meyer
et
al.,
2001
Dualp
hase
ste
el
fcc
cin
tob
cc
a,
here
both
prim
ary
a,
bain
itic
aand
mart
ensite
EB
SD
:in
div
idualO
DF
of
pro
duct
phases;
XR
D:
textu
reof
reta
ined
c
Sta
rtin
gfr
om
various
cte
xtu
res,
thre
eB
ain
variants
,no
variant
sele
ction
Prim
ary
afo
rms
from
defo
rmed
c;b
ain
ite
and
mart
ensite
pre
fera
bly
form
from
recry
sta
llised
cH
utc
hin
son
et
al.,
1998;
Wate
rschoot
et
al.,
2002
Low
carb
on
mic
roallo
yed
ste
el
fcc
cin
top
rim
ary
and
bain
itic
bcc
aE
BS
D:
ind
ivid
ualO
DFs
of
pro
duct
phases;
XR
D:
Sta
rtin
gfr
om
various
cte
xtu
res;
24
KS
variants
;no
variant
sele
ction
Prim
ary
and
bain
itic
afo
rmfr
om
defo
rmed
c;re
tain
ed
cis
part
ially
recry
sta
llised
Mesp
lont
and
De
Coom
an,
2003
0. 3
C–9N
iste
el
bcc
ain
tofc
cc
into
near-
bcc
a’m
art
ensite
EB
SD
:O
DF,
mic
rostr
uctu
reand
mic
rote
xtu
re;
XR
D:
avera
ge
textu
re
Sta
rtin
gfr
om
various
cte
xtu
recom
ponents
;24
KS
variants
,no
variant
sele
ction
Mod
elp
red
ictions
are
consis
tent
with
exp
erim
ent:
no
variant
sele
ction
Yokota
et
al.,
2005
AIS
I304
sta
inle
ss
ste
el
fcc
cin
tohcp
eand
near
bcc
a’m
art
ensite
EB
SD
:lo
cald
efo
rmation
mechanis
ms
(inclu
din
gp
hase
transfo
rmation);
XR
D:
volu
me
fraction
of
phases
and
avera
ge
textu
reof
auste
nite
Sta
rtin
gfr
om
exp
erim
enta
lc
textu
re;
24
variants
ded
uced
from
aP
TM
C,
coup
led
variant
sele
ction
and
cry
sta
lp
lasticity
of
auste
nite,
localis
ation
rule
by
aself-c
onsis
tent
schem
efo
rth
eauste
nite
poly
cry
sta
l
Mod
elp
red
ictions
are
part
ially
consis
tent
with
exp
erim
ent
concern
ing
the
orienta
tion
of
auste
nite
gra
ins
that
transfo
rmfirs
t;d
esp
ite
that
the
mod
elis
not
yet
ab
leto
take
em
art
ensite
into
account
Petit
et
al.,
2006
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
International Materials Reviews 2007 VOL 52 NO 2 113
Coupling EBSD with other advanced techniques:a summaryThe following key issues should be considered in thenext future.
Think in 3D
Not only should experimental techniques be furtherdeveloped, but models should be improved, inasmuch asthey are frequently based on 2D assumptions about, forexample, the morphology of newly formed phases andthe determination of nucleation sites from 2D observa-tions. Crystal connectivity should be necessarily con-sidered from a 3D point of view. Phenomena such asvariant selection should include 3D considerations, todifferentiate ‘true’ variant selection from stereologyartefacts. In situ investigations are still in progress,to be combined with 3D considerations as detailedabove.
Model improvements
Micromechanical modelling of phase transformationsis a favourite frame to use EBSD data combiningcrystallographic and spatial information. However, themost advanced micromechanical models are not oftenused together with EBSD. As an example, PTMCsare often used with approximations on mechanical pro-perties (e.g. stiffness, strength, strain rate and tempera-ture effects) of parent and product phases. They areoften used as if they were descriptive, although theyare phenomenological in nature. On the other hand,localisation and homogenisation rules used in multi-scale modelling of resulting properties could also beimproved.
Concluding remarksAn exponential increase continues in the number ofpublished papers on application of EBSD to phasetransformations, a topic that is still developing rapidly.From ‘revisiting’ a number of features of phasetransformations, EBSD has progressed and now pro-vides original results in a number of specific areas. Asfor other topics related to EBSD, a careful literaturereview is required before starting a new study. The vastmajority of the literature is devoted to metallicmaterials; promising results already exist for ceramicsdespite experimental difficulties. It would be interestingto cross-reference what is commonly found with metallicand ceramic materials.
Electron backscatter diffraction is now a maturetechnique. It is easy to use (except for minerals or for 3Dinvestigations), although handling the amount of dataproduced by 2D or 3D EBSD mapping is increasinglydifficult. Coupling EBSD with in situ investigations,local property measurement and advanced modelling isnow required to get deeper insight into phase transfor-mation mechanisms.
The development of high speed systems and com-puter aided volume reconstruction techniques stronglysuggests that many results on phase transformationcrystallography and mechanisms should be revisitedusing 3D characterisation of partially or fully trans-formed microstructures, which provide input datafor process optimisation and prediction of productproperties.T
ab
le34
Co
nti
nu
ed
Mate
rial
Ph
ase
tran
sfo
rmati
on
Exp
eri
men
tal
data
Mo
del
Resu
lts
Ref.
Ti–
6A
l–4V
bcc
bin
tohcp
aE
BS
Dand
XR
D:
avera
ge
textu
reS
tart
ing
from
bte
xtu
re;
Burg
ers
variants
Ag
reem
ent
with
exp
erim
ent
dep
end
son
bor
(az
b)
initia
lsta
teand
bd
efo
rmation
sta
teb
efo
rep
hase
transfo
rmation
Hum
bert
et
al.,
1994;
Gey
et
al.,
1996;
Mousta
hfid
et
al.,
1997a;
Sta
nfo
rdand
Bate
,2004
Ti
hcp
ain
tob
cc
bE
BS
D:
avera
ge
textu
res
Sta
rtin
gfr
om
ate
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re;
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ers
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ponent
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ere
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ate
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row
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ard
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2004
IMI8
34
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rtin
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stic
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based
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ain
{0001} a
and
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sare
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nific
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ain
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2006
Cu–40Z
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nie
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XR
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ibution
of
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esses
from
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ng
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ht
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zone
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icte
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od
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aet
al.,
1999
Ni–
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into
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aft
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ing
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variants
,111
.b
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reg
ives
the
rig
ht
fib
res
for
the
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reof
c’S
akata
et
al.,
2001
*H
SLA
:hig
hstr
eng
thlo
wallo
y.
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
114 International Materials Reviews 2007 VOL 52 NO 2
Appendix
Orientation relationshipsThe most commonly encountered ORs between phasesin solid state transformations are reported in Table 35.
Acknowledgements
Technical assistance from Mrs Odile Adam for papertracking is gratefully acknowledged. Many thanks aredue to Professor Pineau for his kind and continuoussupport to achieve the present work. The present paperis dedicated to the memory of late Professor Flower,who kindly gave me my first opportunity to makeextensive EBSD measurements in his research group atImperial College.
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Table 35 Commonly encountered ORs in EBSD studies
Name and acronymCrystal structureof one phase
Crystal structure ofthe other phase Parallel planes Parallel directions
Burgers bcc hcp {011}bcc//{0001}hcp ,111.bcc//,1120.hcp
Kurdjumov–Sachs (KS) bcc fcc {011}bcc//{111}fcc ,111.bcc//,110.fcc
Nishiyama–Wassermann (NW) bcc fcc {011}bcc//{111}fcc ,011.bcc//,211.fcc
Pitsch bcc fcc {011}bcc//{001}fcc ,111.bcc//,110.fcc
Greninger–Troiano (GT) bcc fcc {011}bcc ,1u from {111}fcc Midway between KSand closest NW
Nishiyama–Wassermann (NW) fcc hcp {111}fcc//{0001}hcp ,112.fcc//,1100.hcp
Blackburn c (L10) a2 (DO19) {111}c//{0001}a2,110.c//,1120.a2
Pitsch–Petch a (bcc) c (Tetragonal Fe3C) {001}c//{215}a ,100.c 2.6u from ,311.a
and ,010.c 2.6u from,131.a
Bagaryatsky a (bcc) c (Tetragonal Fe3C) {001}c//{112}a ,100.c//,110.a and,010.c//,111.a
Pitsch c (fcc) c (Tetragonal Fe3C) ,100.c//,554.c and,010.c//,010.c and,001.c//,225.c
Gourgues-Lorenzon Application of electron backscatter diffraction to the study of phase transformations
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208. D. Helary, O. Dugne, X. Bourrat, P. H. Jouneau, F. Cellier:
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212. S. Henry, G.-U. Gruen, M. Rappaz: ‘Influence of convection on
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214. D. Hernandez, M. Dıaz-Fuentes, B. Lopez, J. M. Rodriguez-
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215. H. T. Hesemann, P. Mullner, E. Arzt: ‘Stress and texture
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216. R. L. Higginson, B. Roebuck, E. J. Palmiere: ‘Texture develop-
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217. M. Hillert et al.: ‘Preface to the viewpoint set on: bainite’ and
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218. A. Hirose, D. Nakamura, H. Yanagawa, K. F. Kobayashi:
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222. R. A. Holt, P. Zhao: ‘Micro-texture of extruded Zr–2.5Nb tubes’,
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224. K. Horikawa, K. Yoshida: ‘Visualization of hydrogen distribution
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226. Y. Huang, F. J. Humphreys, I. Brough: ‘The application of a hot
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227. M. Humbert, H. Moustahfid, F. Wagner, M. J. Philippe:
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