# Effect of Phosphorus on Intermediate Temperature Embrittlement of High Silicon Spheroidal Graphite Cast Irons

Rodolfo González-Martínez, Jon Sertucha, Yannick Thébault, Jacques Lacaze

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Rodolfo González-Martínez, Jon Sertucha, Yannick Thébault, Jacques Lacaze. Effect of Phosphorus on Intermediate Temperature Embrittlement of High Silicon Spheroidal Graphite Cast Irons. Metal-lurgical and Materials Transactions A, 2025, 10.1007/s11661-025-07804-6 hal-05082752

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# Effect of Phosphorus on Intermediate Temperature Embrittlement of High Silicon Spheroidal Graphite Cast Irons

RODOLFO GONZÁLEZ-MARTÍNEZ, JON SERTUCHA, YANNICK THÉBAUT, and JACQUES LACAZE

Ferritic spheroidal graphite cast irons can be susceptible to intermediate temperature embrittlement between 350 °C and 500 °C. It has been reported that the corresponding loss of ductility increases with the sulfur and magnesium content of the alloy but can be eliminated if the phosphorus content is adapted to ensure a sufficiently low ratio of magnesium to phosphorus contents. Using a low-sulfur, low-phosphorus and high-silicon alloy, previous work clarified the role of magnesium and suggested that small MgO precipitates, developing at grain boundaries due to diffusion of oxygen from the atmosphere during tensile tests, contributed to the generation of brittleness. The present work aimed to clarify how the addition of phosphorus could effect this process and make spheroidal graphite cast irons immune to brittleness at intermediate temperatures. The findings suggest that Mg-P precipitates appeared at high temperature during the cooling process after casting, and that they catch excess magnesium at grain boundaries. During mechanical testing, these Mg-P precipitates tighten the oxygen that penetrates along the grain boundaries to form large Mg-P-O precipitates that are not detrimental to ductility, while preventing the massive formation of small Mg-O precipitates.

https://doi.org/10.1007/s11661-025-07804-6
© The Author(s) 2025

## I. INTRODUCTION

Low temperature embrittlement and intermediate temperature embrittlement are of definite concern for many applications of spheroidal graphite irons (SGI). Low temperature embrittlement is mostly characterized by the transition from ductile to brittle (TDB) temperature below which ductility falls to zero and impact toughness achieves quite low values. Unfortunately, the TDB temperature for ferritic SGIs increases from about – 100 °C at 1 wt pct Si to about room temperature at 2.7 wt pct Si [1] This low temperature embrittlement is much alike that of ferritic steels, and phosphorus in SGIs has the same detrimental effect of increasing the TDB temperature. [2–4]

Intermediate temperature embrittlement (ITB) is a loss of ductility characterized by a marked reduction in elongation at rupture in a tensile test. It occurs at temperatures between 350 °C and 500 °C, with an upward shift as the strain rate increases. Wright and Farrell[5] pointed out that not all ferritic SGIs are sensitive to ITB, as was subsequently emphasized on several occasions and indicated in our previous work.[6] ITB can be appropriately defined by an elongation at rupture being lower than a critical value of about 10 pct, and it is characterized by a large amount of intergranular fracture (IG) instead of the expected ductile fracture at those temperatures. High-silicon and SiMo SGIs are also potentially sensitive to ITB, and this led to a wealth of works on its impact on thermal fatigue that was recently reviewed by Lekakh et al.[7]

Wright and Farrell[5] pointed out that ITB is associated with high magnesium and high sulfur contents, and several works reported the presence of Mg-bearing precipitates on the fracture surface[8–10] and on the intergranular facets.[6,11,12] These precipitates have been identified by transmission electron microscopy (TEM) as MgO precipitates.[11] Some authors proposed that cracks initiated on voids where these precipitates formed upon casting[8] while others suggested that magnesium segregates at the ferrite grain boundaries and weakens them.[10] A major step was achieved in the 1990s when it

<footnote>RODOLFO GONZÁLEZ-MARTÍNEZ and JON SERTUCHA are with the AZTERLAN, Basque Research and Technology Alliance (BRTA), Durango, Spain YANNICK THÉBAUT and JACQUES LACAZE are with the CIRIMAT, Université de Toulouse, BP 44362, 31030 Toulouse, France. Contact e-mail: jacques.lacaze@toulouse-inp.fr
Manuscript submitted December 2, 2024; accepted May 2, 2025.</footnote>
was discovered that addition of phosphorus at more
than 0.032 wt pct[¹³,¹⁴] or better such as pct Mg/pct
P < 1.5[¹⁵⁻¹⁷] makes SGIs immune to ITB (where pct Mg
and pct P are the content of the alloy in magnesium and
phosphorus expressed in weight percent). This critical
ratio was later found to be around 1.0.[¹⁸] In these
modified alloys, P-containing precipitates have been
observed,[¹⁸,¹⁹] as well as Mg-rich precipitates when the
critical ratio is exceeded.[¹⁸] Several authors[¹⁸⁻²⁰] sum-
marized ITB by stating that magnesium poisons ferrite
grain boundaries, but that this detrimental effect can be
counteracted by addition of phosphorus that is prone to
tight magnesium in compounds. It is quite interesting to
notice that other elements of the same column as
phosphorus can act in the same way, namely arsenic
(Japanese patent by Takanezawa et al. mentioned by Li
and Sloss[²¹]) and antimony.[¹⁹] However, arsenic is very
hazardous and should not be present in modern cast
irons, while antimony added at the necessary level is a
powerful pearlite former whose addition should be
limited when a ferritic matrix is sought in the as-cast
state.

The above analysis based on magnesium segregation to
the ferrite grain boundaries disregards the result by
Wright and Farrell that a SGI prone to ITB under air is
immune to it under vacuum. This discrepancy was
highlighted in our previous work[⁶] and we suggested that
it was the MgO precipitates that caused the intergranular
fracture. Further, as the temperature was increased for
tests under air, it was observed that these precipitates
became more numerous and enlarged to reach about
200 nm in size at 450 °C. They are very much similar in
size and distribution to the precipitates that can be
observed on figure 10 of Wright and Farrell[⁵] on the
rupture surface of an alloy tensile tested at 370 °C. Part of
this figure is reproduced in Figure 1(a) where it can be
compared with our previous observations for a test at

450 °C (Figure 1(b)). It was therefore concluded that ITB
failure results from MgO precipitates spread over the
ferrite grain boundaries and growing due to diffusion of
oxygen from the outside atmosphere during testing, until
reaching a critical size where ITB occurs. The sensitivity
of ITB to strain rate mentioned above is consistent with
such a diffusion-controlled mechanism.

As reminded above, ITB can be suppressed if enough
phosphorus has been added to the alloy. This appears at
first quite striking because phosphorus is known to be
detrimental for low temperature toughness. Bak *et al.*[22]
investigated this on a SGI with high phosphorus level at
0.07 wt pct and evidenced grain boundary segregation of
this element by holding the material for 100 hours at
450 °C. This has been further studied with Auger
analysis by Lejcek *et al.*[23] who concluded that phos-
phorus segregation initiates at high temperature in the
austenitic domain and is enhanced as the temperature
decreases. However, such intergranular segregation is a
sluggish process needing long holding times that are one
to two orders of magnitude longer than the time needed
for tensile testing and may not interfere with the
mechanism of suppressing ITB. The present work was
intended to clarify this issue by analyzing how addition
of phosphorus can immune SGIs to ITB, and this was
done by tensile testing SGIs with various amounts of
added phosphorus.

II. EXPERIMENTAL DETAILS

Three melt baches of about 60 kg were successively
prepared following the same experimental procedure,
materials, devices and tools as it was previously
described[⁶] though two different additions of FeP alloy
(grain size 8 to 30 mm, P = 23.8, Mn = 2.22, Ti =
2.21, Si = 1.36, Mg = 0.65, V = 0.28, Ca = 0.18,

<figure><img src="image_2.png" /><figcaption>Fig. 1—IG facets covered with precipitates: (a) TEM replica after figure 10 from Wright and Farrell[5] on an alloy tensile tested at 370 °C; (b) SEM micrograph of a high-silicon SGI tested at 450 °C during previous work.[6] The micrograph in (a) was originally published in the “1985 *AFS Transactions*” (Paper #85-137) and is reprinted with permission from the American Foundry Society, Inc, Schaumburg, Illinois (www.afsin.c.org).</figcaption></figure>
<table><caption>Table I. Chemical Composition (Wt Pct) and Pct Mg/Pct P Ratio of the Three Alloys Prepared in This Work</caption><thead><tr><th>Alloy</th><th>C</th><th>Si</th><th>Mn</th><th>P</th><th>S</th><th>Mg</th><th>Cr</th></tr></thead><tbody><tr><td>1P</td><td>3.13</td><td>3.59</td><td>0.13</td><td>0.014</td><td>0.008</td><td>0.041</td><td>0.039</td></tr><tr><td>2P</td><td>3.01</td><td>3.80</td><td>0.14</td><td>0.027</td><td>0.006</td><td>0.043</td><td>0.034</td></tr><tr><td>3P</td><td>3.12</td><td>3.57</td><td>0.15</td><td>0.042</td><td>0.007</td><td>0.041</td><td>0.031</td></tr></tbody></table><table><thead><tr><th>Alloy</th><th>Ni</th><th>Cu</th><th>Ti</th><th>Ce</th><th>La</th><th>Sb</th><th>Pct Mg/Pct P</th></tr></thead><tbody><tr><td>1P</td><td>0.020</td><td>0.023</td><td>0.034</td><td>0.0084</td><td>0.0032</td><td>0.0055</td><td>2.93</td></tr><tr><td>2P</td><td>&lt; 0.020</td><td>0.024</td><td>0.025</td><td>0.0085</td><td>0.0031</td><td>0.0076</td><td>1.59</td></tr><tr><td>3P</td><td>&lt; 0.020</td><td>0.021</td><td>0.036</td><td>0.0080</td><td>0.0031</td><td>0.0081</td><td>0.98</td></tr></tbody></table>

Cr = 0.16, Ni = 0.15, Cu = 0.15 and Fe balance, wt pct) were made to obtain the alloy compositions shown in Table I. The amounts of FeP added were 49.3 g and 109.9 g for alloys 2P and 3P, respectively (no addition of FeP was made for alloy 1P). These additions were made to the metallic charge initially introduced into the melting furnace and the P content was then checked with the rest of the elements at the end of the melting stage. In addition, 2.5 to 2.6 g of Sb was added to the melting furnace for all three batches to prevent degeneration of the shape of the graphite spheroids.[24]

All prepared melts were treated with a FeSiMg alloy (grain size 5 to 25 mm, Si = 46.60, Mg = 6.00, Ca = 0.96, Al = 0.71 and rare earth RE = 0.92, Fe balance, wt pct) following the sandwich method and by transferring 50 kg of the prepared alloy to a ladle. Before that, the spheroidizing alloy was placed in the chamber of the ladle in an amount of 0.60 kg (about 1.2 wt pct of the treated batch weight) and then covered using steel scrap pieces 5 to 15 mm in size. Once the spheroidizing reactions were finished, the melts were rapidly cleaned and cast in seven chemically bonded sand molds to produce Y2 wedges according to the EN-1563 standard. Each of these molds contained one cavity in which 14 g (0.20 wt pct of the total weight of the alloy poured in the mold) of a commercial inoculant (grain size 0.2 to 0.5 mm, Si = 69.9, Al = 0.93, Ca = 1.38, Bi = 0.49, RE = 0.37 and Fe balance, wt pct) was added before pouring.

After cooling of the wedges, they were removed from the molds and cleaned and the microstructure at room temperature was checked with light optical metallography. Two tensile specimens (total length of 115 mm, gage of 50 mm in length and 10 mm in diameter) were machined out from the bottom part of each casting which is the zone with the lowest risk of inclusions and shrinkage porosities. Tensile tests to determine ultimate tensile strength (UTS), yield strength (YS) and elongation at rupture (A) at 50 °C, 100 °C, 350 °C, 450 °C, 550 °C and 650 °C were performed on these specimens after a 10 min holding at temperature. The lowest temperature of 50 °C was selected instead of room temperature (RT) to avoid making experiments in between the lower and upper shelves of the brittle to ductile low temperature transition. The tensile parameters UTS, YS and A were measured using an MTS 810 tensile testing equipment at a constant displacement rate of 0.90 mm/min, thus corresponding to an initial strain rate of $3 \times 10^{-4} s^{-1}$. For temperature testing, a MTS653 furnace attached to the tensile equipment was used. Three thermocouples located along the

gage section were used to monitor the temperature of the specimens and to limit the temperature fluctuations within ± 3 °C.

Once tested, one of the fracture surfaces of each specimen was analyzed by scanning electron microscopy (SEM) using a Quanta 200 Carl Zeiss Ultraplus operated at 20 kV that has a field emission gun (FEG). Chemical microanalyses were carried out with an energy dispersive spectrometer (EDS) X-max 20 Oxford attached to the SEM equipment. The other part of the tensile specimens were cut to obtain metallographic samples at about 5 to 7 mm away from the rupture surface. After polishing, a representative picture at 100× magnification (0.757 mm² each) per sample was taken without etching for image analysis of graphite particles using the ImageJ software. For this, all graphite particles with an area equal to or higher than 25 µm² were counted, then assigned to classes III, V or VI according to the criteria described with figure 7 in the work by Germain *et al.*[25] Nodularity values (pct) were determined using the expression $Nod = \frac{N_v+N_{vi}}{N_{III}+N_v+N_{vi}} \times 100$, where $N_{III}$, $N_v$ and $N_{vi}$ are the number of particles assigned to classes III, V and VI, respectively. On the other hand, nodule count (mm⁻²) values were calculated as $N = \frac{N_v+N_{vi}}{0.757}$. These metallographic samples were finally etched using Nital 5 pct reactant to check the amounts of ferrite and pearlite.

Further SEM observations have been carried out on longitudinal sections of the ruptured samples P1-350 and P3-350 close to the fracture surface. Use was made of a FEI Quanta 450 microscope equipped with a Quantax 30 mm² SDD EDS detector from Bruker that was operated at 20 kV. Another system combined a JEOL 7800F Prime FEG-SEM with an Oxford Instruments AZtecEnergy SDD X-Max 80 mm² EDS detector, which was operated at 5 kV to reduce the interaction of the beam with the material and thus increase EDS resolution. In all cases, EDS spectra were accumulated until a total count between 100,000 and 200,000 was obtained. Depending on the system, the Y axis of the spectra can be either the total count or the number of counts per eV per second.

### III. RESULTS

#### A. *Effect of Temperature on the Mechanical Properties*

In the as-cast state, all alloys had a fully ferritic microstructure, i.e., a matrix with less than 5 pct
pearlite, as illustrated with the micrograph of alloy 2P in Figure 2. For all three alloys, the distribution of graphite particles showed Nod and N values in the ranges 91 to 96 pct and 220 to 245 mm⁻², respectively, indicating that the successive steps of melt preparation, spheroidizing treatment and inoculation were satisfactorily processed.

All tensile curves recorded on alloys 1P, 2P and 3P at the different selected temperatures are reported in Appendix I, including a plot of the whole curves and a zoom on the elastic part. Table II lists the characteristic tensile values that were obtained. The stress-strain curves recorded with alloys 1P and 3P are compared in Figure 3. For alloy 3P (dashed curves), the ultimate tensile strength changed little up to 350 °C and then decreases with temperature, while for alloy 1P (solid curves) it decreases monotonously with temperature. This continuous decrease for alloy 1P is in relation with the marked dip in elongation at rupture at less than 5 pct at 350 °C (see Table II), while this parameter remains at more than 15 pct for alloy 3P at that temperature. These results therefore support the role of phosphorus in the elimination of ITB. Alloy 2P, with intermediate phosphorus content, showed an evolution with temperature which confirms the positive effect of this element, with a value of elongation at rupture at 350 °C that was intermediate at about 9 pct.

<figure><img src="image_3.png" /><figcaption>Fig. 2—Light optical micrograph of the microstructure of alloy 2P after Nital etching.</figcaption></figure>

Values of A in Table II are also plotted in Figure 4(a). For the three alloys of the present study, the ductility loss appears limited to around 350 °C, while it extended up to 500 °C for the very high-silicon alloy (4.9 wt pct) that has been previously investigated under the same conditions.[6] Figure 4(b) shows the UTS values for the three alloys and it is quite noticeable that the effect of ITB shows up at 350 °C, with the UTS values ranked as the A values. In our previous work,[6] it was emphasized that most of the literature results that were analyzed showed little variation of UTS up to 350 °C as noticed here for alloy 3P. The reason for this behavior becomes evident with the present results when noticing the jerky undulations of some strain-stress curves in Figure 3. These undulations are seen for all three alloys at 200 °C, and also at 350 °C for alloy 3P. They are due to flow instabilities characteristic of dynamic strain ageing (DSA) that has been described long ago for mild steels (see for example the review by Leslie[26]) and recently re-analyzed by Caillard[27] who used in-situ observations of dislocation movements during straining and clarified the effect of alloying elements in Fe-C alloys (e.g., silicon, nickel, chromium).

While low temperature DSA of cast irons has been recognized as such for a long time (see the review by Jalava et al.[28]), jerky flow was not seen in our previous work and was not reported in the analyzed studies,[6]

<figure><img src="image_4.png" /><figcaption>Fig. 3—Effect of temperature on the stress-strain curves of alloys 1P (solid lines) and 3P (dashed lines).</figcaption></figure>

<table><caption>Table II. Tensile Test Results (UTS, YS and A) Conducted at 50 °C, 200 °C, 350 °C, 450 °C, 550 °C and 650 °C for the Three 1P, 2P and 3P Alloys</caption><thead><tr><th rowspan="2">Temperature (°C)</th><th colspan="3">UTS (MPa)</th><th colspan="3">YS (MPa)</th><th colspan="3">A (Pct)</th></tr><tr><th>1P</th><th>2P</th><th>3P</th><th>1P</th><th>2P</th><th>3P</th><th>1P</th><th>2P</th><th>3P</th></tr></thead><tbody><tr><td>50</td><td>517</td><td>541</td><td>521</td><td>391</td><td>423</td><td>401</td><td>24.6</td><td>23.1</td><td>23.4</td></tr><tr><td>200</td><td>467</td><td>504</td><td>509</td><td>335</td><td>367</td><td>365</td><td>17.9</td><td>17.5</td><td>14.8</td></tr><tr><td>350</td><td>384</td><td>475</td><td>536</td><td>313</td><td>354</td><td>376</td><td>4.3</td><td>8.9</td><td>15.2</td></tr><tr><td>450</td><td>271</td><td>332</td><td>338</td><td>268</td><td>327</td><td>313</td><td>12.1</td><td>19.2</td><td>18.7</td></tr><tr><td>550</td><td>183</td><td>154</td><td>169</td><td>169</td><td>153</td><td>155</td><td>22.9</td><td>18.8</td><td>27.9</td></tr><tr><td>650</td><td>56</td><td>71</td><td>105</td><td>54</td><td>67</td><td>99</td><td>15.4</td><td>26.7</td><td>25.6</td></tr></tbody></table>
<figure><img src="image_5.png" /><figcaption>Fig. 4—Evolution with temperature of the elongation at rupture (a) and of the ultimate tensile strength (b) of the investigated alloys.</figcaption></figure>

even though when serrations were clearly present[19] or
probable.[29] Appearance of mechanical flow instabilities
depends on the strain rate as well as on the thermo-me-
chanical history of the material, which might explain
why it is not systematically observed. Extending the
literature review to studies devoted to low-cycle
thermo-mechanical fatigue or other service properties
showed numerous examples.[30–33] The detailed study of
Chen et al.[30] showed mechanical flow instabilities occur
in cast irons as in mild steels, being controlled by carbon
diffusion to the dislocation core at low temperature
(200 °C) while depending on pipe diffusion of substitu-
tional solutes at higher temperature (350 °C). Cail-
lard[27] clarified the role of these substitutional solutes
depending on the fact that they have attractive or
repulsive interaction energy with carbon. In complex
alloys as are cast irons, occurrence or absence of flow
instabilities at high temperature (300 °C to 400 °C) may
also depend on the relative amounts of low-level
elements (e.g., S, P, Mn, Cu) and their interactions in
the process of pipe diffusion.

B. Characterization of the Precipitates in the As-Cast Materials

From the results reported above, the three alloys
prepared in this work are suitable for studying the
positive effect of phosphorus on ITB. Initial observa-
tions were made by looking for precipitates in the
as-cast materials, as shown in Figure 5. Alloy 1P
(Figure 5(a)) contains small precipitates of MgO oxide,
as well as (MgAlSi)N nitrides and some TiC carbides.
Characterization of the precipitates was carried out
using EDS and Figure 5(d) shows an example of an EDS
spectrum recorded on a particle identified as MgO with
P, Sb and rare earth elements (RE) also present at very
low levels. In this spectrum as in all those that follow,
peaks of C, Si and Fe appear which are attributed to the
surrounding matrix.

In addition to the TiC particles, the 2P intermediate
alloy (Figure 5(b)) contains both MgO precipitates and
precipitates containing significant amounts of Mg, O
and P, which we will refer to as Mg-O-P. Figure 5(e)
shows a typical EDS spectrum of these precipitates,
which also contain the elements Sb and RE. Finally, the
same precipitates were also found in alloy 3P with
maximum sizes comparatively larger than those detected
in alloy 2P (see Figure 5(c)). It is relevant to indicate that
quite small Mg-O-P particles were found here and there
in alloy 1P (not shown in Figure 5(a)) and also in the
previously studied alloy with 0.012 wt pct P.[6] In
summary, the interesting feature in these as-cast mate-
rials is that the maximum size of the P-containing
precipitates increases as P content does, i.e., from alloy
1P to alloy 3P.

C. Description of the Rupture Surfaces

Figure 6 shows the rupture surfaces of alloys 1P, 2P and 3P that have been tested at the lowest studied temperature of 50 °C and at 350 °C. At 50 °C (column to the left, a to c), mostly ductile fracture features were obtained with primary cupules having a size in direct relation with the graphite spheroids. The ligaments were also found to contain secondary micro-cupules that appear to have developed from precipitates in the matrix. The rupture surfaces were scanned looking for IG facets even though they could not be numerous owing to the high values of the elongation at rupture. Accordingly, some rare IG facets could be found in alloy 1P and very few in alloy 2P such as the one pointed out with a black arrow in Figure 6(b), while no IG facet could be seen on alloy 3P ruptured at 50 °C.

On the other hand, a continuous evolution of the
fracture surface was observed for the tests at 350 °C
(right-hand column in Figure 6), from alloy 1P, which
mainly exhibited IG facets (Figure 6(d)), to alloy 3P, for
which the fracture remained essentially ductile
(Figure 6(f)). For this temperature, precipitates were
<figure><img src="image_6.png" /><figcaption>Fig. 5—SEM micrographs obtained from alloy 1P (a), alloy 2P (b) and alloy 3P (c). Note that the scale is the same for all three images. Typical EDS record for MgO (d) and Mg-O-P (e) particles.</figcaption></figure>

easily observed on the IG facets (alloys 1P and 2P), at the grain boundaries and in the ligaments of the ductile zones.

Figure 7 illustrates, at higher magnification, the IG facets observed on the 1P and 2P alloys tested at 50 °C. They show small, scattered precipitates that closely resemble those previously reported and analyzed as MgO precipitates, see Figure 1. As in our previous work,[6] the Mg and O signals in the EDS spectra (not

shown here) were quite weak due to the small size of these precipitates, but sufficiently clear to support this conclusion.

As expected from the low values of the elongation at rupture, the rupture surface of alloys 1P and 2P tested at 350 °C and 450 °C showed many IG facets with MgO precipitates. In full agreement with previous observations,[6] the MgO precipitates grew in number and size with the testing temperature. This is demonstrated by
<figure><img src="image_7.png" /><figcaption>(a)</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>20 µm</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>(b)</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>20 µm</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>(c)</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>20 µm</figcaption></figure>

<caption>(f)</caption>

Fig. 6—SEM micrograph (secondary electrons) of the rupture surface of alloy 1P (a and d), alloy 2P (b and e) and alloy 3P (c and f) tested at 50 °C (a, b and c) and at 350 °C (d, e and f).

comparing the micrographs of sample 2P tested at 450 °C (Figure 8(a)) and 50 °C (Figure 7(b)) considering that their scales are slightly different. While the rupture surface was essentially ductile in alloy 3P with the highest P content as mentioned above, some IG facets could also be found on which some MgO precipitates were detected. This is illustrated in Figure 8(b) for the sample tested at 450 °C.

Appendices 2 and 3 show further SEM micrographs of the fracture surface at 350 °C and 450 °C, respectively, with focus on the observation of IG facets at higher enlargement. Finally, the rupture surfaces were fully ductile for all three alloys tested at 550 °C (not shown).

Attention was then directed to the comparatively large precipitates present in the ligaments, and it was noticed they could have size up to 10 µm and also
<figure><img src="image_13.png" /><figcaption>Fig. 7—In-lens SEM micrographs of an IG facet present in alloy 1P (a) and 2P (b) tensile ruptured at 50 °C.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>Fig. 8—MgO precipitates present on IG facets of alloy 2P (a) and 3P (b) tested at 450 °C.</figcaption></figure>

sometimes appear clearly multiphasic. Figure 9(a) illustrates a few of them related to micro-cupules present in alloy 3P tested at 50 °C. The spectrum #35 was acquired on the flank of a cupule to serve as a reference for the matrix with only C, Fe and Si detected. The spectrum #32 showed mostly Mg and O peaks and was very similar to that in Figure 5(d), with faint traces of Ti, P and Sb. Spectrum #34 was quite alike that in Figure 5(e), with marked Mg and O peaks but also significant levels of P and Sb, as well as Ce. It was noted that, in that case, the O peak was significantly larger than the Mg peak. Finally, spectrum #33 was the most complicated (see Figure 9(b)) as it contains the Mg and O peaks and also significant P, Sb, Ce and La peaks, and also a strong Ti peak suggesting the presence of TiC in addition to Mg and P-bearing phases. Again, the O peak was larger than the Mg peak in the spectrum #34.

Similarly to what was made for the samples tested at 50 °C, focus was then put on the large particles found in the samples tested at 350 °C, such as the ones shown in Figures 10(a) and (b) in the case of alloy 2P. The EDS

spectrum of all of the particles that were observed shows the presence of Mg, O, P and Ce, with also Sb and La in some cases. Also, Ti and other carbide former elements were noticed in some cases, associated with a stronger C peak suggesting additional precipitation of a carbide. The elemental maps in Figure 10(c) recorded on the particle shown in Figure 10(b) suggest that this precipitate was in fact multiphasic while it was not always as clear from other similar precipitates.

From all the above observations, the following findings could be proposed:

- When only intense Mg and O peaks appear on the EDS spectrum in addition to minor Fe, Si and C peaks, being the peak for Mg and O of similar height, the analysed precipitate should be MgO oxide.

- When a Ti peak appears, this is due to the presence of a TiC precipitate (with possibly other carbide former elements such as V, Mo, and Cr), and the carbon signal should accordingly be high. In many
<figure><img src="image_15.png" /><figcaption>Fig. 9—Secondary electron SEM micrograph of alloy 3P ruptured at 50 °C with location of various EDS measurements (see text) (a) and EDS spectrum at location #33 (b).</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>Fig. 10—Black arrows indicate large particles found on a IG facet in alloy 2P tested at 350 °C (a) and enlarged view of one of these particles (b). Elemental maps of O, Mg, P and Ce recorded on the particle shown in (b) (c).</figcaption></figure>

cases, TiC particles appeared associated with Mg–O–P precipitates that might be due to rejection of Ti from this latter, or else to rejection of Mg and P from the former.

- P is always present with Mg, and Sb, Ce and/or La are also present to some extent.

Emphasis has been put on the relative height of O and Mg peaks and it was noticed that the Mg peak becomes much lower than the O peak when P is present to a

significant level. It should be appreciated that this correlation was used as an indication because the relative heights of the peaks actually depend on all the other elements present in the area analysed. With this limitation in mind, the above observation suggests that the presence of Mg, O and P, with the O peak significantly larger than the Mg peak, is related to a biphasic precipitation of MgO and an Mg–O–P phase. This latter could certainly be based on Mg₃P₂ that has been reported by Yan *et al.*,¹⁸ possibly being a
<figure><img src="image_17.png" /><figcaption>Fig. 11—Secondary crack in 1P sample ruptured at 350 °C.</figcaption></figure>

phosphate such as the Mg₃(PO₄)₂. Similar observations were made on alloys 1P and 2P but the relative distribution of MgO and P-bearing precipitates could not be evaluated while very few of these latter would be expected in alloy 1P.

D. *Observations Close to the Rupture Surface of Materials Tested at 350 °C*

Longitudinal diametrical sections of samples 1P and 3P tested at 350 °C were prepared aiming at characterizing the area close to the rupture surface where penetration of oxygen is expected to have occurred. In 1P sample, a few secondary short cracks could be observed slightly away from the rupture surface, while no distributed damage could be seen in 3P sample that ruptured along a well-defined and unique crack. An example of a secondary crack in sample 1P is seen in Figure 11. Along this crack, IG facets were easily

<figure><img src="image_18.png" /><figcaption>Fig. 12—(a, and b) SEM micrographs of precipitate alignments close to a secondary crack in sample 1P tested at 350 °C. The dashed arrows indicate locations with composition as that of the matrix but enriched in oxygen. (c) EDS spectrum of a MgO precipitate in (a) and (d) EDS spectrum of the complex precipitate of MgO + Mg₃P₂ in (b).</figcaption></figure>
<figure><img src="image_19.png" /><figcaption>Fig. 13—(a) Precipitates in sample 3P tested at 350 °C. (b) Higher magnification of one of the arrowed MgO precipitates in (a).</figcaption></figure>

<figure><img src="image_20.png" /><figcaption>Fig. 14—(a) Higher magnification SEM micrograph of the Mg-O-P precipitate in Figure 13(a). (b) Example of the EDS spectra recorded at the locations indicated with plus sign in (a).</figcaption></figure>

identified that often showed the same features as in Figure 7(a). However, a close examination evidenced that all MgO precipitates had spalled off during sample preparation which might indicate that their bonds to the matrix were not strong.

Attention was then turned to the surroundings of secondary cracks such as those two seen in Figure 12 that both present alignment of precipitates certainly related to grain boundaries. EDS spectra were acquired on many of these precipitates that showed in some cases alignment of MgO precipitates (Figure 12(a)) or evidenced a variety of precipitates (Figure 12(b)). At several locations in this 1P sample, areas with a grey contrast similar to that of the matrix had a rougher, more irregular surface such as those indicated with a dashed arrow in the micrographs of Figure 12. EDS spectra of such areas showed similar to that of the

matrix but with a marked oxygen peak evidencing that oxygen penetration had occurred along grain boundaries during mechanical testing. It can be noticed that the MgO precipitates that appear in dark grey contrast in Figure 12(a) are separated by light grey areas with similar contrast to that for the matrix enriched in oxygen, and this was indeed verified by EDS, whilst this was not seen in Figure 12(b) apart close to the cavity.

In the area of Figure 12(a), P was found as a Mg₃P₂ precipitate surrounded by a TiC precipitate away from the alignment of MgO precipitates, and there was no oxygen detected there. This two-phase precipitate had certainly formed early during the casting process and got engulfed in austenite during solidification. Strikingly, the EDS spectra of the MgO precipitates in the alignment in Figure 12(a) did not show any P or any other element than Mg and O and those of the matrix
(Fe, C and Si); see the example in Figure 12(c). That there was no P in these MgO precipitates strongly suggests that the grain boundary marked with these precipitates appeared during the eutectoid transformation as a boundary between two grains of ferrite growing inside an austenite grain. In contradistinction, the presence of P in the precipitates of Figure 12(b) indicates they are located along a grain boundary between ex-austenite grains. Furthermore, that the matrix along the grain boundary between the P-bearing precipitates in Figure 12(b) did not show oxygen enrichment evidences an effect of phosphorus on oxygen distribution in such boundaries.

Alignments of precipitates were not so easily found in the 3P sample tested at 350 °C. Figure 13(a) is an example of an area with several precipitates that could well feature the presence of grain boundaries. EDS spectra were recorded and the particles have been labelled accordingly. P was not recorded in the upper right MgO precipitate but it was in the two others that are arrowed. At higher magnification, it could be noticed the presence of bright spots (Figure 13(b)) indicating that the precipitates grew over another phase that was probably Mg₃P₂.

The elongated Mg-O-P precipitate in the center of Figure 13(a) is seen at higher magnification in Figure 14(a) and strongly resembles the precipitate shown in Figure 10(b). Though it seems multiphasic, the three spectra recorded at the locations indicated with a plus sign were all similar, and an example is given in Figure 14(b) where the O peak appears again far larger than the Mg peak. Note that the bottom recording was expected to give a different spectrum, due to the brighter appearance of the precipitate there; the fact that this is not the case may be due to a slight shift of the beam during acquisition.

Using the internal standard of the Oxford EDS, a semi-quantitative analysis was obtained that gave the following atomic ratio for the spectrum in Figure 14(b) of 63 pct O, 17 pct Mg and 10 pct P. This result is not far from the atomic ratio in the Mg₃(PO₄)₂ phase that is 62 pct O, 23 pct Mg and 15 pct P and thus confirms the nature of this phase. We can therefore assume that the oxygen diffusing along the grain boundaries was preferentially trapped by the Mg₃P₂ precipitates to transform to Mg₃(PO₄)₂ phosphide. A partial transformation may have occurred during cooling after casting (Figure 5(e)), when the testing temperature was low (Figure 9(b)) or when the oxygen content was too low (Figure 12(d)).

#### IV. DISCUSSION

The discussion below will focus on the suppression of ITB by the addition of phosphorus to form precipitates of Mg-O-P. There is, however, a prerequisite for this discussion, which is to clarify the impact of DSA on the loss of ductility and on ITB occurrence. The review in the recent study by Jalava et al.[28] does not provide a definitive answer in the case of low-cycle fatigue, as DSA can be beneficial or detrimental, though both

<figure><img src="image_21.png" /><figcaption>Fig. 15—Effect of strain rate on the temperatures for ITB onset, ductility minimum and recovery (drawn after Wright and Farrell[5]). Symbols have been added that represent the upper temperature for serrated flow as reported in the indicated works.</figcaption></figure>

ferritic and ferrito-pearlitic matrices were considered to arrive to this conclusion. If only ferritic alloys under static tensile tests are considered, it is useful to remind that the ITB domain shifts to higher temperature when the strain rate is increased. This is illustrated in Figure 15, adapted from Wright and Farrell,[5] where the intergranular fracture onset limit, ductility minimum and recovery limit are shown. On the same graph have also been reported with symbols the maximum temperature at which flow instabilities were observed in various reports.[30,34,35] It is seen that there exists a close correlation between serrated flow and onset of ITB, though it should also be realized that the ductility decrease associated with ITB may not have allowed for flow instabilities to set up during tensile test. In their pioneering work, Yanagisawa and Lui[35] arrived to the conclusion that the onset of ITB is related to both a micro-mechanical process—build-up of stresses around the graphite spheroids that is enhanced by DSA- and a metallurgical process. However, the fact that some SGIs present ITB without DSA, while others show DSA and no ITB, suggests that the metallurgical process is predominant in the appearance of ITB. Following our previous work that established the role of intergranular precipitation of MgO on the occurrence of ITB, the discussion below elaborates on the mechanism by which addition of phosphorus alleviates ITB.

All elements entering in the Mg-bearing precipitates that have been reported above, namely Mg, O and P, but also Ce, La and Sb, do segregate positively during austenite solidification, thus get enriched in the last zones to solidify. These elements are likely to combine to form compounds, either during solidification or during solid-state cooling. Therefore, these precipitates will be preferentially observed in the ex-austenite grain boundaries at room temperature after the eutectoid transformation. As an example, Lekakh et al.[19] studied a 4 pct
Si eutectic cast iron doped with either P or Sb and reported Mg-P and Mg-Sb precipitates in the corresponding alloy in the as-cast state. The authors also noticed the frequent presence of Ce and La in the EDS analyses, as was also the case in the present observations.

Because of the known ability of RE to form com-
pounds with other elements, we searched the literature
for additional information in relation to ITB. Specific
results on Ce alone were found in two reports, one by
Chen et al.[11] and the other by Li and Sloss.[21] Chen
et al.[11] investigated ITB sensitivity of cast irons
spheroidized with pure Mg, pure Ce, and two with a
mix of them. They reported disappearance of ITB in the
material that had been spheroidized with Ce only while
the other three alloys showed a strong ductility decrease
with a minimum at 400 °C. The authors concluded that
even the mere presence of 0.02 wt pct Mg was sufficient
to generate IG fracture that is characteristic of ITB. Li
and Sloss[21] investigated the role of Ce at 0.033 wt pct
added and concluded that it behaves as Mg in triggering
ITB. However, their result was obtained on an alloy also
containing 0.018 wt pct Mg and does not affect the
conclusion by Chen et al. In the alloys that contained
both Ce and Mg, Chen et al. noticed the presence of
both MgO and CeP and do not mention Mg-P particles,
which suggests that Ce first associates most of the
phosphorus as CeP particles making it unavailable to
tight magnesium that instead forms detrimental MgO
precipitates. The finding by Chen et al. does not appear
to have been disseminated, probably because the nodu-
larity achieved with pure Ce treatment is insufficient.

The above indicates that in industrial practice where
the RE content is quite low, some P could be caught at
high temperature, but enough will remain to interfere with
Mg and O. As an example of practical interest, the study
by Lin et al.[10] is worth mentioning. These authors carried
out cyclic heating of a 4.0 wt pct Si SGI (with 0.045 wt pct
P and 0.060 wt pct Mg) and noted the presence of both
MgO and of Ce-bearing Mg-O-P particles. After typi-
cally 20 cycles up to an upper temperature set in between
923 K (650 °C) and 1123 K (850 °C), the samples were
tensile tested at RT. The authors reported a strong
decrease of the RT elongation at rupture for an upper
temperature of 1023 K (750 °C) associated with full IG
rupture. EDX-SEM analysis evidenced MgO and
Mg-O-P particles and their Figure 8 shows that the
crack path goes through grain boundaries filled with MgO
particles and not along those containing Mg-O-P pre-
cipitates. This is a strong support to the findings shown in
the present work.

In summary, it can be suggested that addition of P to
a SGI alloy leads to the early formation of Mg₃P₂
precipitates, and possibly also of MgO and Mg-O-P
precipitates (see Figure 5) in the last to freeze areas, i.e.,
at the outer austenite grain boundaries. During tensile
testing, oxygen will penetrate along grain boundaries
and lead to further precipitation of MgO, either as small
discrete particles in low-P alloys or mostly on top on the
existing Mg-O-P precipitates on high-P alloys. In the
former case, MgO precipitates will preferentially appear
along the grain boundaries generated during the eutec-
toid transformation that will be low in phosphorus and

this precipitation will lead to ITB. In the latter case,
oxygen is preferentially bound by Mg₃P₂ precipitates to
give large phosphides that do not appear to lead to ITB.
Provided the amount of cerium is low enough, addition
of phosphorus to mitigate excess magnesium is success-
ful to make SGI immune to ITB but should be limited as
this element leads to low temperature embrittlement.

V. CONCLUSION

A group of three alloys containing increasing P
content has been studied to investigate the expected
beneficial effect of this element to reduce the interme-
diate temperature embrittlement (ITB) of high silicon
ferritic ductile irons. In agreement with literature, it was
observed that the ductility loss decreases with the ratio
pct Mg/pct P present in the alloy, that is, when
increasing the P content. SEM analysis of the metal
matrix near the fracture surface shows that MgO
particles are present with TiC in all three alloys,
although significant amounts of large Mg-O-P precip-
itates were only found in the 2P and 3P alloys. A
comparison of these Mg-O-P particles indicate that
their maximum size increases with the P content. The
inspections performed on the fracture surfaces at 350 °C
and 450 °C revealed the presence of relevant amounts of
MgO particles on the IG facets of alloys 1P and 2P
which size increased with the testing temperature. For
alloy 3P at the same testing temperature only some IG
facets were found with some MgO precipitates. MgO
and large Mg-O-P particles were found in the rupture
ligaments.

It is finally suggested that the addition de P to a SGI
alloy leads to the early formation of Mg₃P₂ precipitates,
and possibly also of MgO and Mg-O-P precipitates in
the last solidification areas. During tensile testing under
air, oxygen will penetrate along grain boundaries and
lead to further precipitation of MgO, either as small
discrete particles in low-P alloys or mostly on top on the
existing Mg-O-P precipitates on high-P alloys. In the
former case, MgO precipitates will preferentially appear
along the grain boundaries generated during the eutec-
toid transformation that will be low in phosphorus and
this precipitation will lead to ITB. In the latter case,
oxygen is preferentially bound by Mg₃P₂ precipitates to
give large phosphides that appear to reduce ITB.

ACKNOWLEDGMENTS

Alexandre Freulon of CIRIMAT is warmly thanked
for the metallographic preparation of the longitudinal
section of the samples tested at 350 °C. The valuable
comments of the key reader and reviewers were highly
appreciated.

CONFLICT OF INTEREST

On behalf of all authors, the corresponding author
states that there is no conflict of interest.
APPENDIX 1: STRAIN-STRESS CURVES
OF THE THREE INVESTIGATED ALLOYS

See Figure 16.

<figure><img src="image_22.png" /><figcaption>Fig. 16—Whole stress-strain curves and zoom on the elastic part for alloys 1P (a and b), 2P, (c and d) and 3P (e and f).</figcaption></figure>
APPENDIX 2: RUPTURE SURFACES
OF THE SAMPLES TESTED AT 350 °C

See Figure 17.

<figure><img src="image_23.png" /><figcaption>Fig. 17—SEM micrograph (secondary electrons) of the rupture surface of alloy 1P (a and d), alloy 2P (b and e) and alloy 3P (c and f) tested at 350 °C.</figcaption></figure>
APPENDIX 3: RUPTURE SURFACES
OF THE SAMPLES TESTED AT 450 °C

See Figure 18.

<figure><img src="image_24.png" /><figcaption>Fig. 18—SEM micrograph (secondary electrons) of the rupture surface of alloy 1P (a and d), alloy 2P (b and e) and alloy 3P (c and f) tested at 450 °C.</figcaption></figure>
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